Composite materials comprising two jonal functions and methods for making the same

ABSTRACT

The present invention generally relates to mechanisms for preventing undesirable oxidation (i.e., oxidation protection mechanisms) in composite bodies. The oxidation protection mechanisms include getterer materials which are added to the composite body which gather or scavenge undesirable oxidants which may enter the composite body. The getterer materials may be placed into at least a portion of the composite body such that any undesirable oxidant approaching, for example, a fiber reinforcement, would be scavenged by (e.g., reacted with) the getterer. The getterer materials) may form at least one compound which acts as a passivation layer, and/or is able to move by bulk transport (e.g., by viscous flow as a glassy material) to a crack, and sealing the crack, thereby further enhancing the oxidation protection of the composite body. One or more ceramic filler materials which serve as reinforcements may have a plurality of super-imposed coatings thereon, at least one of which coatings may function as or contain an oxidation protection mechanism. Specifically, a coating comprising boron nitride which has been engineered or modified to contain some silicon exhibits improved corrosion resistance, specifically to oxygen and moisture. The coated materials may be useful as reinforcing materials in high performance composites to provide improved mechanical properties such as fracture toughness. The present invention also relates to improved composites which incorporate these materials, and to their methods of manufacture.

CROSS REFERENCE TO RELATED APPLICATIONS

The present patent disclosure is a continuation-in-part of U.S. patentApplication Ser. No. 08/472,613, filed on Jun. 7, 1995, now U.S. Pat.No. 5,682,594, which issued on Oct. 28, 1997, in the names ofChristopher R. Kennedy et al., and entitled “Composite Materials andMethods For Making The Same”.

This invention was made with Government support under Contract No.DE-FC02-92CE40994 awarded by the Department of Energy. The Governmenthas certain rights to this invention.

TECHNICAL FIELD

The present invention generally relates to filler materials which areadapted for use as the reinforcement phases(s) in composite bodies.Coated ceramic filler materials comprised of ceramic particles, fibers,whiskers, etc. having at least two substantially continuous coatingsthereon are provided. The coatings are selected so that the interfacialshear strength between the ceramic filler material and the firstcoating, between coatings, or between the outer coating and thesurrounding matrix material, are not equal so as to permit debonding andpull-out when fracture occurs. The resultant, multicoated ceramic fillermaterials may be employed to provide composites, especially ceramicmatrix composites with increased fracture toughness. The ceramic fillermaterials are designed to be particularly compatible with ceramicmatrices formed by directed oxidation of precursor metals, but suchceramic filler materials are also adaptable for use in many othercomposite material systems.

The present invention also relates to techniques for increasing thecorrosion resistance of composite materials, particularly of ceramicfiber reinforced composites exposed to oxygen and water vapor atelevated temperatures. One approach to inhibiting corrosion in a ceramicmatrix composite body is to reduce the number and/or size of microcracksin the body, thereby reducing access of corrodants to the interior ofthe body. Another broad approach is to provide chemical additives to thebody which are capable of gettering a corrodant or interfering with itscorrosion mechanisms.

BACKGROUND ART

A ceramic composite is a heterogeneous material or article comprising aceramic matrix and filler such as ceramic particles, fibers or whiskers,which are intimately combined to achieve desired properties. Thesecomposites are produced by such conventional methods as hot pressing,cold pressing and firing, hot isostatic pressing, and the like. However,these composites typically do not exhibit a sufficiently high fracturetoughness to allow for use in very high stress environments such asthose encountered by gas turbine engine blades.

A novel and useful method for producing self-supporting ceramiccomposites by the directed oxidation of a molten precursor metal isdisclosed in Commonly Owned U.S. Pat. No. 4,851,375, which issued onJul. 25, 1989, described below in greater detail. However, theprocessing environment is relatively severe, and there is a need,therefore, to protect certain fillers from the strong oxidationenvironment. Also, certain fillers may be reduced at least partially bymolten metal, and therefore, it may be desirable to protect the fillerfrom this local reducing environment. Further, the protective meansshould be conducive to the metal oxidation process, yet not degrade theproperties of the resulting composite, and even more desirably provideenhancement to the properties. Still further, in some instances it maybe desirable for the means or mechanisms for protecting the fillerduring matrix or composite formation to also protect the fillers againstundesirable attack of oxidants diffusing through the matrix duringactual service of the composite.

It is known in the art that certain types of ceramic fillers serve asreinforcing materials for ceramic composites, and the selection orchoice of fillers can influence the mechanical properties of thecomposite. For example, the fracture toughness of the composite can beincreased by incorporating certain high strength filler materials, suchas fibers or whiskers, into the ceramic matrix. When a fractureinitiates in the matrix, the filler at least partially debonds from thematrix and spans the fracture, thereby resisting or impeding theprogress of the fracture through the matrix. Upon the application ofadditional stress, the fracture propagates through the matrix, and thefiller begins to fracture in a plane different from that of the matrix,pulling out of the matrix and absorbing energy in the process. Pull-outis believed to increase certain mechanical properties such aswork-of-fracture by releasing the stored elastic strain energy in acontrolled manner through friction generated between the material andthe surrounding matrix.

Debonding and pull-out have been achieved in the prior art by applying asuitable coating to the ceramic filler material. The coating is selectedso as to have a lower bonding strength with the surrounding matrix thanthe filler, per se, would have with the matrix. For example, a boronnitride coating on silicon carbide fibers has been found to be useful toenhance pull-out of the fibers. Representative boron nitride coatings onfibers are disclosed in U.S. Pat. No. 4,642,271, which issued on Feb.10, 1987, in the name of Roy W. Rice, and are further disclosed in U.S.Pat. No. 5,026,604, which issued on Jun. 25, 1991, in the name ofJacques Thebault. However, the use of boron nitride coated fibers incomposites may present significant processing disadvantages. Forexample, the production of ceramic matrix composites containing boronnitride coated materials requires the use of reducing atmospheres sincea thin layer of boron nitride readily oxidizes (e.g., converts to boronoxide in an oxygen-containing atmosphere) at temperatures above 800-900°C. A reducing atmosphere, however, may often times not be compatiblewith the directed oxidation of molten parent metal for fabricatingceramic composites. Further, in the directed oxidation process thecoating desirably is compatible with the molten metal in that the moltenmetal wets the coated filler under the process conditions, for otherwisethe oxidation process and matrix growth may be impeded by the filler.

Another drawback of boron nitride is that, upon oxidation, the boriareaction product can dissolve or further react with water to form boricacid, which can be a vapor under the local oxidizing conditions. Thus,the boria is not a passive layer, but can be continually removed throughvolatilization. U.S. Pat. No. 5,593,728 to Moore et al. addresses thisshortcoming of boron nitride. Specifically, by producing a pyrolytic BNcoating containing from 2 to 42 wt % silicon, with substantially no freesilicon present, Moore et al. observe greatly reduced rates of oxidativeweight loss. The coating is formed by CVD using reactant vapors ofammonia and a gaseous source of both boron and silicon. The gases areflowed into a reaction chamber between a temperature of 1300° C. and1750° C. and within a pressure range of 0.1 Torr to 1.5 Torr.

It is not clear, however, whether the modified BN layer of Moore et al.permits molten parent metal to wet the coating (for infiltration) andyet resist any adverse reaction therewith. Further, the modified BNcoatings of Moore et al. were deposited onto single filaments. Due tothe high deposition rates resulting from the deposition conditions, itis unclear whether the Moore et al. technique could be applied to coat aplurality of fibers, e.g.a stack of fabrics making up a preform.

Also, in order to prevent or minimize filler degradation, certain limitsmay be imposed on the conventional fabrication processes, such as usinglow processing temperatures or short times at processing temperature.For example, certain fillers may react with the matrix of the compositeabove a certain temperature. Coatings have been utilized to overcomedegradation, but as explained above, the coating can limit the choice ofprocessing conditions. In addition, the coating should be compatiblewith the filler and with the ceramic matrix.

A need therefore exists to provide coated ceramic filler materials whichare capable of debonding and pull-out from a surrounding ceramic matrix.A further need exists to provide coated ceramic filler materials whichmay be incorporated into the ceramic matrix at elevated temperaturesunder oxidizing conditions to provide composites exhibiting improvedmechanical properties such as increased fracture toughness.

In order to meet one or more of these needs, the prior art shows fillermaterials bearing one or more coatings. Carbon is a useful reinforcingfiller but typically is reactive with the matrix material. It thereforeis well known in the art to provide the carbon fibers with a protectivecoating. U.S. Pat. No. 4,397,901, which issued on Aug. 9, 1983, in thename of James W. Warren, teaches first coating carbon fibers with carbonas by chemical vapor deposition, and then with a reaction-formed coatingof a metallic carbide, oxide, or nitride. Due to a mismatch in thermalexpansion between the fiber and the coating, the fiber is capable ofmoving relative to the coating to relieve stress. A duplex coating oncarbon fibers is taught by U.S. Pat. No. 4,405,685, which issued on Sep.20, 1983, in the names of Honjo et al. The coating comprises a first orinner coating of a mixture of carbon and a metal carbide and then anouter coating of a metal carbide. The outer coatings prevent degradationof the fiber due to reaction of unprotected fiber with the matrixmaterial, and the inner coating inhibits the propagation of cracksinitiated in the outer layer. U.S. Pat. No. 3,811,920, which issued onMay 21, 1974, in the names of Galasso et al. relating to metal matrixcomposites, discloses coated fibers as a reinforcing filler, such asboron filaments having a silicon carbide surface layer and an additionalouter coating of titanium carbide. This reference teaches that theadditional coating of titanium carbide improves oxidation resistance aswell as provides a diffusion barrier between the filament and metalmatrix.

However, the prior art fails to teach or suggest filler materials with aduplex coating for protection from potentially corrosive environmentsduring manufacture or operation of the composite body and yet in thecomposite material permit the filler to debond and pull-out from thesurrounding matrix. Moreover, the prior art does not recognize certainother oxidation protection mechanisms which can be employed jointly.Specifically, the prior art fails to appreciate certain importantaspects of utilizing getterer materials which function to scavengeundesirable oxidants, and optionally after such scavenging has occurred,forming desirable compounds or materials (e.g., one or more glassycompounds) which assist in protecting the reinforcement materials fromundesirable oxidation.

DESCRIPTION OF COMMONLY OWNED U.S. PATENTS AND PATENT APPLICATIONS

The filler materials utilized in this invention may be protected by anumber of different mechanisms in a number of different compositebodies. Filler materials containing a coating or plurality of coatings,in accordance with the teachings of this invention, are particularlyapplicable or useful in the production of ceramic composites disclosedand claimed in Commonly Owned U.S. Pat. No. 4,851,375, entitled “Methodsof Making Composite Ceramic Articles Having Embedded Filler,” whichissued on Jul. 25, 1989, from U.S. patent application Ser. No. 819,397,filed Jan. 17, 1986, which is a continuation-in-part of Ser. No.697,876, filed Feb. 4, 1985 (now abandoned), both in the names of MarcS. Newkirk et al. and entitled “Composite Ceramic Articles and Methodsof Making Same”. This Commonly Owned Patent discloses a novel method forproducing a self-supporting ceramic composite by growing an oxidationreaction product from a precursor metal or parent metal into a permeablemass of filler.

The method of growing a ceramic product by an oxidation reaction of aparent metal is disclosed generically in Commonly Owned U.S. Pat. No.4,713,360, which issued on Dec. 15, 1987, in the names of Marc S.Newkirk et al. and entitled “Novel Ceramic Materials and Methods ofMaking Same”; and in U.S. Pat. No. 4,853,352, which issued on Aug. 1,1989, in the names of Marc S. Newkirk et al. and entitled “Methods ofMaking Self-Supporting Ceramic Materials”.

Commonly Owned U.S. Pat. No. 4,713,360 discloses a novel method forproducing a self-supporting ceramic body by oxidation of a parent metal(as defined below) to form an oxidation reaction product which thencomprises the ceramic body. More specifically, the parent metal isheated to an elevated temperature above its melting point but below themelting point of the oxidation reaction product in order to form a bodyof molten parent metal which reacts upon contact with a vapor-phaseoxidant to form an oxidation reaction product. The oxidation reactionproduct, or at least a portion thereof which is in contact with andextends between the body of molten parent metal and the oxidant, ismaintained at the elevated temperature, and molten metal is drawnthrough the polycrystalline oxidation reaction product and towards theoxidant, and the transported molten metal forms oxidation reactionproduct upon contact with the oxidant. As the process continues,additional metal is transported through the polycrystalline oxidationreaction product formation thereby continually “growing” a ceramicstructure of interconnected crystallites. Usually, the resulting ceramicbody will contain therein inclusions of nonoxidized constituents of theparent metal drawn through the polycrystalline material and solidifiedtherein as the ceramic body cooled after termination of the growthprocess. As explained in these Commonly Owned Patents and PatentApplications, resultant novel ceramic materials are produced by theoxidation reaction between a parent metal and a vapor phase oxidant,i.e., a vaporized or normally gaseous material, which provides anoxidizing atmosphere. In the case of an oxide as the oxidation reactionproduct, oxygen or gas mixtures containing oxygen (including air) aresuitable oxidants, with air usually being preferred for obvious reasonsof economy. However, oxidation is used in its broad sense in theCommonly Owned Patents and in this application, and refers to the lossor sharing of electrons by a metal to an oxidant which may be one ormore elements and/or compounds. Accordingly, elements other than oxygenmay serve as the oxidant. In certain cases, the parent metal may requirethe presence of one or more dopants in order to influence favorably orto facilitate growth of the ceramic body, and the dopants are providedas alloying constituents of the parent metal. For example, in the caseof aluminum as the parent metal and air as the oxidant, dopants such asmagnesium and silicon, to name but two of a larger class of dopantmaterials, are alloyed with the aluminum alloy utilized as the parentmetal.

The aforementioned Commonly Owned U.S. Pat. No. 4,853,352 discloses afurther development based on the discovery that appropriate growthconditions as described above, for parent metals requiring dopants, canbe induced by externally applying one or more dopant materials to thesurface or surfaces of the parent metal, thus avoiding the necessity ofalloying the parent metal with dopant materials, e.g. metals such asmagnesium, zinc and silicon, in the case where aluminum is the parentmetal and air is the oxidant. External application of a layer of dopantmaterial permits locally inducing metal transport through the oxidationreaction product and resulting ceramic growth from the parent metalsurface or portions thereof which are selectively doped. This discoveryoffers a number of advantages, including the advantage that ceramicgrowth can be achieved in one or more selected areas of the parentmetal's surface rather than indiscriminately, thereby making the processmore efficiently applied, for example, to the growth of the ceramicplates by doping only one surface or only portions of a surface of aparent metal plate. This improvement invention also offers the advantageof being able to cause or promote oxidation reaction product growth inparent metals without the necessity of alloying the dopant material intothe parent metal, thereby rendering the process feasible, for example,for application to commercially available metals and alloys whichotherwise would not contain or have appropriately doped compositions.

In forming a ceramic composite body, as described in the aforesaidCommonly Owned U.S. Pat. No. 4,851,375, the parent metal is placedadjacent a permeable mass of filler material, and the developingoxidation reaction product infiltrates the mass of filler material inthe direction and towards the oxidant and boundary of the mass. Theresult of this phenomenon is the progressive development of aninterconnected ceramic matrix, optionally containing some nonoxidizedparent metal constituents distributed throughout the growth structure,and an embedded filler.

In producing the ceramic composite, any suitable oxidant may beemployed, whether solid, liquid, or gaseous, or a combination thereof.If a gas or vapor oxidant, i.e. a vapor-phase oxidant, is used thefiller is permeable to the vapor-phase oxidant so that upon exposure ofthe bed of filler to the oxidant, the gas permeates the bed of filler tocontact the molten parent metal therein. When a solid or liquid oxidantis used, it is usually dispersed through a portion of the bed of filleradjacent the parent metal or through the entire bed, typically in theform of particulates admixed with the filler or as coatings on thefiller particles.

Polycrystalline bodies comprising a metal boride are produced inaccordance with Commonly Owned U.S. Pat. No. 4,777,014, which issued onOct. 11, 1988, in the names of Marc S. Newkirk, et al., and entitled“Process for Preparing Self-Supporting Bodies and Products MadeThereby”. In accordance with this invention, boron or a reducible metalboride is admixed with a suitable inert filler material, and the moltenparent metal infiltrates and reacts with the boron source. This reactiveinfiltration process produces a boride-containing composite, and therelative amounts of reactants and process conditions may be altered orcontrolled to yield a polycrystalline body containing varying volumepercents of ceramic, metal, reinforcing filler, and/or porosity.

U.S. Pat. No. 5,202,059 to Kennedy et al. teaches ceramic fillermaterials having a plurality of superimposed coatings thereon. Thecoated materials are useful as reinforcing materials in ceramic matrixcomposites to provide improved mechanical properties such as fracturetoughness. The coatings are selected so that the interfacial shearstrength between the ceramic filler material and the first coating,between coatings, or between the outer coating and the surroundingmatrix material, are not equal so as to permit debonding and pull-outwhen fracture occurs. By reason of this invention, the coated ceramicfiller materials not only provide improved mechanical properties, butalso the filler is protected from severe oxidizing environments and yetamenable to the process conditions for the manufacture of the ceramiccomposite.

The entire disclosures of each of the Commonly Owned Patents and PatentApplications are incorporated herein by reference.

SUMMARY OF THE INVENTION

In accordance with the present invention, there is disclosed a pluralityof distinct, but combinable, mechanisms for preventing undesirableoxidation (i.e., oxidation protection mechanisms) of reinforcementmaterials (e.g., fibers) in composite bodies. These oxidation protectionmechanisms include the use of getterer materials which are present in atleast a portion of the composite body (e.g., in at least a portion ofthe matrix; in at least a portion of one or more interfacial coatings;or in, on or adjacent to at least a portion of the reinforcingmaterials, etc.). These getterer materials tend to scavenge (e.g., reactwith) undesirable oxidants which enter the composite body such as, forexample, by diffusion mechanisms, through microcracks, etc. Theseoxidation protection mechanisms may, in certain embodiments, alsoinclude techniques for reducing the number and/or size of suchmicrocracks in a portion of or throughout the composite body. Thereduction in the amount and/or size of microcracks limits the transportof undesirable oxidants into and out of the composite body.

When a composite body is put into service in an oxidizing environment,and assuming that the oxidizing environment would have an adverse effectupon the reinforcing material, some type of oxidation protectionmechanism should be utilized to prevent the reinforcement from oxidizingundesirably. If a getterer material was placed on, or at least in closeproximity to, the reinforcing material, then an oxidant which came intocontact with the getterer material, such as by diffusion mechanisms,through microcracks, etc., could be gettered (e.g., reacted) by thegetterer materials, thereby ameliorating undesirable reaction(s) withthe reinforcing material. Further, if the getterer material forms acompound, such as for example, a glass, the compound could provide evenfurther oxidation protection to the reinforcing material. In thisregard, if a glass so formed had an appropriate viscosity, then theformed glass could flow into any microcracks which may be present nearthe glass, thereby permitting the glass to function as a crack sealant.Such desirable compounds are often termed “trap sealants”. In thisregard, the formed glass ideally has an oxidant permeability which issufficiently low to provide for suitable oxidation protection at theintended operation temperatures of the composite body for the desiredamount of time.

In another embodiment of the invention, a glassy material or aglass-forming material is provided to the composite body duringfabrication.

The composite body can be engineered so that one or more getterermaterials are included in the composite body such that one or moredesirable compounds (e.g., glasses) are formed. Each of the getterermaterials could react with one or more oxidants at differenttemperatures and form one or more desirable compounds (e.g., one or moredesirable glasses) which may provide for differing amounts of oxidationprotection at different temperatures. In addition, the formed compoundscould further react with other species contained in the composite bodyto produce additional desirable compounds. Further, such a formedcompound could react or interact (e.g., alloy) with a glass orglass-former material which may have been provided to the composite bodyduring fabrication. Accordingly, a composite body could be producedwhich contained a plurality of different oxidation protectionmechanisms, wherein each oxidation protection mechanism was included toprovide for desirable oxidation protection at different servicetemperatures of the composite body.

One exemplary manner of placing an oxidant getterer material onto areinforcing material would be to dip, paint or spray an appropriatematerial onto at least a portion of the reinforcing material prior tomatrix formation. Alternatively, chemical vapor deposition (CVD) orchemical vapor infiltration (CVI) techniques could be utilized to obtainone or more coatings on at least a portion of, or in a preferredembodiment, substantially all of, a reinforcing material. It would bedesirable for such coatings to be capable of surviving any matrixformation steps in addition to providing in-service oxidationprotection. Moreover, such coating could contain, or be modified tocontain as, for example, one or more additional species. Such speciesmight function as an oxygen getterer or may provide oxidation resistanceby some other mechanism.

In a preferred embodiment of the invention, a coated ceramic fillermaterial, adaptable for use as a reinforcing component in a ceramicmatrix or metal matrix composite, is provided with a plurality ofsuperimposed coatings. The filler or reinforcing material useful forthis embodiment includes materials where the length exceeds thediameter, typically in a ratio of at least about 2:1 and more preferablyat least about 3:1, and includes such filler materials as whiskers,fibers, and staple. The coating system includes a first coating insubstantially continuous contact with the ceramic filler material, andone or more additional or outer coatings superimposed over theunderlying coating, and in substantially continuous contact therewith.Zonal junctions are formed between the filler and first coating, betweensuperimposed coatings, and between the outer coating and the ceramicmatrix. The coatings are selected so that the interfacial shear strengthof at least one of these several zones is weak relative to the otherzones. As used herein and in the appended claims, a zonal junction isnot limited to an interface, per se, between the surfaces but alsoincludes regions of the coatings in proximity to the interfaces, andshear, therefore, is zonal in that it may occur at an interface orwithin a coating. Further, it is understood that the zonal junctionbetween adjacent surfaces may be minimal or negligible and exhibitessentially no bonding or adhesion, or the adjacent surfaces may exhibitappreciable bonding or a strong bond. Upon the application of fracturestress to the composite, the weak zone allows for debonding of thefiller before the filler fractures, and pull-out or shear of the fillerupon fracture of the filler. This debonding and friction pull-outenhances certain mechanical properties of the composite, and inparticular debonding improves the fracture toughness. Thus, in a duplexcoating system, for example, having a first coating and a second, outercoating superimposed on the first coating, the coatings are chosen tofacilitate debonding and pull-out such that junction between one of thethree interfaces (i.e. the interface between the filler and the innercoating, the interface between the inner coating and the outer coating,the interface between the outer coating and the surrounding matrix, orthe strength of a coating) is weak relative to the other zonal junctionsand allows for debonding and pull-out.

By reason of this embodiment of the invention, the coated ceramic fillermaterials not only provide improved mechanical properties, but also thefiller is protected from severe corrosive environments during use andyet amenable to the processing conditions for making a composite (e.g.,matrix formation). For example, in developing a ceramic matrix bydirected metal oxidation, certain fillers and/or coatings thereon may beat least partially reduced by the molten parent metal upon contact, andthus the outer coating protects the filler and inner coating againstthis local reducing environment. Thus, duplex coated fillers areadaptable for use as a reinforcing component in a ceramic matrixcomposite formed by the directed oxidation reaction of a moltenprecursor metal or parent metal with an oxidant. For many of the samereasons, such duplex coated filler materials are adaptable for use inmetal matrix composite systems in which the metallic matrix is formed byinfiltration.

Coated fillers also find utility in composite materials formed by aninfiltration process where the filler material is not wetted by theinfiltrant. In composite systems featuring the melt infiltration ofsilicon based metals, the infiltrating silicon alloy will wet siliconcarbide fillers, for example, but does not readily wet other usefulfillers such as ceramic oxides. Moreover, duplex (or higher order)coated fillers may find utility in composite systems where the matrix isformed by an infiltration process but where an inner coating on thefiller (provided, for example, for de-bonding the filler from thematrix) may not be wetted by the infiltrant material. Boron nitride, forexample, makes a desirable debond coating, but boron nitride is notreadily wet by metals such as aluminum or silicon.

In another preferred embodiment of the invention, the coatings mayprotect the fibers by a means different from, but possibly in additionto, the above-described mechanisms. Specifically, under in-serviceconditions (e.g., at elevated temperatures), the coatings may help topreserve the “original” or “as-fabricated” strength of the fibers bypreserving the original character of the fibers, in particular, thefiber chemistry and/or crystal structural (or lack thereof). Withoutwishing to be bound by any particular theory or explanation, it ispossible that the fiber coating serves to prevent or at least retardthermal decomposition, specifically by preventing, or at leastretarding, outgassing from the fiber, which outgassing may, in somecircumstances, be accompanied by a change in the character of thecrystals making up the fiber such as, for example, through growth ofcertain of the crystals or, in the case of an originally amorphousfiber, by crystallization or devitrification of this amorphousstructure. Specifically, from the perspective of concentrationgradients, a coating on a fiber containing the same elemental species asthe fiber might be expected to retard diffusion of that species out ofthe fiber. For example, a carbon doped boron nitride coating on asilicon carbide based fiber also containing some oxygen and nitrogenmight be expected to reduce the diffusional loss of carbon and nitrogenfrom the fiber.

In general, coated filler materials of this invention may be utilized inthe manufacture of composite materials (e.g., ceramic matrix composites)that provide improved mechanical properties, especially increasedfracture toughness. When so employed, the thickness of the coatingsshould be sufficient to protect the ceramic filler material againstcorrosive environments such as those of molten metals. However, thecoatings should not be so thick as to hinder matrix formation or tointerfere with the function of the filler.

When relatively thick preforms are to be coated by means of CVI, it canbe a challenge sometimes to adequately coat the filler in the center ofthe preform without closing off the pore space between bodies of fillermaterial residing toward the preform exterior (e.g., “canning”), therebyrendering the preform impermeable. Where a preform comprises anassemblage of units, it has been discovered that arranging the unitssuch that the more porous, higher permeability units are situated closerto the preform exterior, and likewise the less porous, lowerpermeability units being situated closer to the center of the preformprovides for a more uniform deposit (thickness-wise) throughout thepreform of reaction product from the reactant gases . Thus, if a preformis to consist of a plurality of woven fabric plies superimposed on topof one another, it would be desirable from the CVI coating uniformityperspective to place the fabric plies having the more “open” weaves onthe outside of those fabric plies having a tighter, less permeableweave.

It is noted that particular emphasis is herein placed upon matricesformed by the directed oxidation of a molten metal, however, certainaspects of the coating composition and/or coating thickness may betransferable to other matrices (e.g., glass matrices, etc.) and/or othermatrix formation conditions (e.g., melt infiltration,chemical vaporinfiltration, etc.).

Moreover, the coatings can be selected so that one or more of thecoatings themselves serves as an oxidant getterer when the composite isput into service. In a further preferred embodiment, once the oxidantgetterer has formed a compound (e.g., at least one glassy compound) dueto a reaction between the getterer and the oxidant, the formed compoundprovides further protection due to, for example, flowing into a crack tofunction as a crack sealant. Still further, the formed compound mayinteract with (e.g., react, alloy or modify) a glass to form a differentglass which could then provided oxidation protection in a differenttemperature regime.

In yet another achievement of the invention, an approach for reducingthe number and/or size of microcracks formed during composite formationand/or formed during composite service or use, is discussed. Microcracksmay be undesirable because such microcracks may permit ready access ofundesirable oxidants to the reinforcement materials which can result indegradation of some properties of the composite body. Specifically,microcracking of a matrix material located between adjacent plies offiber tows/bundles (e.g., silicon carbide fibers) can be reduced orpossibly even eliminated by introducing into the matrix one or morematerials having a relatively low coefficient of thermal expansion suchas, for example, silicon carbide particulate. Thus, to practice thisembodiment of the invention, an appropriate material or combination ofmaterials could be inserted between one or more fiber tows or betweenfiber layers to form a preform from a combination of fibers andparticulate. After formation of the preform, a ceramic matrixcomprising, for example, an oxidation reaction product, could be formed.

The composite bodies of the present invention do not require a seal coatapplied over the exterior of the bulk body. Accordingly, the compositebodies of the present invention are adaptable to finishing operationssuch as machining, polishing, grinding, etc. The resultant compositesare intended to include, without limitation, industrial, structural, andtechnical ceramic bodies for applications where improved strength,toughness and wear resistance are important or beneficial.

While this disclosure focuses primarily on ceramic matrix compositebodies having a matrix formed by the directed oxidation of a moltenmetal, it should be understood that the coating techniques of theinvention are by themselves novel and useful and have industrialapplicability in many other composite body formation processes (e.g.,other ceramic matrix composite formation techniques, glass matrixformation techniques, polymer matrix formation techniques, metal matrixformation techniques, etc.). Accordingly, this invention also relates tothe specific techniques for forming such coatings.

DEFINITIONS

The following terms, as used herein and in the claims, have the statedmeanings as defined below:

The term “oxidation reaction product” in conjunction with both oxidationreaction product growth and gettering means one or more metals in anyoxidized state wherein the metal(s) have given up electrons to or sharedelectrons with another element, compound, or combination thereof.Accordingly, an “oxidation reaction product” under this definitionincludes the product of the reaction of one or more metals (e.g., aparent metal comprising aluminum, silicon, tin, titanium, zirconium,etc.) with an oxidant such as oxygen or air, nitrogen, a halogen,sulfur, phosphorous, arsenic, carbon, boron, selenium, tellurium;compounds such as silica (as a source of oxygen), and methane, ethane,propane, acetylene, ethylene, and propylene (as a source of carbon); andmixtures such as H₂/H₂O and CO/CO₂ which are useful in reducing theoxygen activity of the environment.

The term “oxidant” means one or more suitable electron acceptors orelectron sharers and may be a solid, liquid, or gas (vapor) or somecombination of these. Thus, oxygen (including air) is a suitablevapor-phase gaseous oxidant for the formation of oxidation reactionproduct, with air being preferred for reasons of economy. Boron, boroncarbide and carbon are examples of solid oxidants for the formation ofoxidation reaction product under this definition.

The term “parent metal” refers to that metal, e.g. aluminum, which isthe precursor of a polycrystalline oxidation reaction product such asalumina, and includes that metal or a relatively pure metal, acommercially available metal having impurities and/or alloyingconstituents therein, and an alloy in which that metal precursor is themajor constituent; and when a specified metal is mentioned as the parentmetal, e.g. aluminum, the metal identified should be read with thisdefinition in mind unless indicated otherwise by the context.

The term “ceramic” is not limited to a ceramic body in the classicalsense, that is, in the sense that it consists entirely of non-metallic,inorganic materials, but rather, it refers to a body which ispredominantly ceramic with respect to either composition or dominantproperties, although the body may contain substantial amounts of one ormore metallic constituents such as derived from the parent metal, mosttypically within a range of from about 1-40% by volume, but may includestill more metal.

The term “glass” or “glassy compound” as used in this disclosure broadlyrefers to inorganic materials exhibiting only short-range order, e.g.,non-crystalline character. The term thus includes the traditional“glass-forming oxides” such as silica and boria, but also includesmaterials which normally exhibit long-range order, e.g., crystallinity,but which order has been disrupted through rapid solidification or thepresence of defects such as impurity atoms.

The term “inorganic polymer” or “preceramic polymer” refers to thatclass of polymeric materials which upon pyrolysis convert to ceramicmaterials. Such polymers may be solid or liquid at ambient temperature.Examples of these polymers include the polysilazanes andpolycarbosilazanes which can be pyrolyzed to yield ceramic materialscomprising silicon nitride and silicon carbide, respectively.

The term “melt infiltration” refers to a technique for producingcomposite materials by infiltration whereby a molten metal comprisingsilicon is placed into contact with a permeable mass which is wettableby the molten metal, and the molten metal infiltrates the mass withoutthe requirement for the application of pressure or vacuum. Theinfiltration may occur with significant reaction or with substantiallyno reaction of infiltrating metal and one or more components of thepermeable mass.

The term “reaction-formed” in the context of silicon carbide refers to amelt infiltration process whereby silicon in the infiltrant metal reactswith a carbon source in the permeable mass to produce silicon carbide inthe matrix phase of the resulting composite body. This in-situ formedsilicon carbide may or may not be interconnected.

The term “trap sealant”, as used herein, refers to a chemical specieswhich is capable of gettering oxygen, and upon so doing, forms orcontributes to oxide glass formation, which glass is capable ofinterfering with oxygen gas transport.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a scanning electron micrograph taken at about 350×magnification of a coated ceramic filler material in a ceramic matrixand made according to the invention.

FIG. 2 is a scanning electron micrograph taken at about 850×magnification of ceramic matrix composite having a coated NICALON®ceramic fiber as filler material and made according to the Examplebelow.

FIG. 3 is a scanning electron micrograph taken at 250× magnification ofa fractured surface of the composite made with the coated fibersaccording to the Example below showing extensive pull-out of the fibers.

FIG. 4 is a scanning electron micrograph taken at 800× magnification ofa fractured surface of the composite made with uncoated fibers accordingto the Example below showing no pull-out of the fibers.

FIG. 5a is a schematic of the top view of harness satin weave fabric inthe as-is position as discussed in Example 2.

FIG. 5b is a schematic cross-sectional representation of a harness satinweave fabric in the as-is position as discussed in Example 2.

FIG. 5c is an isometric schematic view illustrating the axes of rotationfor a harness satin weave fabric in the as-is position as discussed inExample 2.

FIG. 5d is a schematic cross-sectional representation of a fabricpreform comprised of harness satin fabric as discussed in Example 2.

FIG. 5e is an isometric schematic representation of a graphitecontainment fixture for effecting the coating of a fabric preform asdiscussed in Example 2.

FIG. 5f is a isometric schematic representation of a cantilever graphitefixture for holding a boron nitride coated fabric preform to enablecoating of the preform with a second coating as discussed in Example 2.

FIG. 5g is a schematic cross-sectional representation of a growth lay-upfor forming a fiber reinforced ceramic composite body as discussed inExample 2.

FIG. 5h is a schematic cross-sectional representation of a lay-up forremoving the metallic component of the formed fiber reinforced ceramiccomposite body discussed in Example 2.

FIG. 6 is a schematic cross-sectional representation of a typical lay-upfor removing at least one metallic constituent of a metallic componentfrom substantially all surfaces of a composite body.

FIG. 7 is an orthoscopic view of tensile and stress rupture testspecimens.

FIG. 8 is a typical stress-strain curve for a fiber-reinforced ceramiccomposite tensile test specimen.

FIG. 9 is a SEM photograph at about 5× magnification of the fracturesurface of a tensile test specimen.

FIG. 10 shows tensile strength of a fiber-reinforced ceramic matrixcomposite vs. T(° C.) in air.

FIG. 11 shows tensile strength vs. temperature for thermally cycled andnon-thermally cycled fiber ceramic matrix composite test specimens.

FIG. 12 shows results of stress rupture testing of NICALON® fiberreinforced Al₂O₃ at 1000, 1100 and 1200° C. in air.

FIG. 13 is a SEM photograph at about 50× magnification of the fracturesurface of a stress rupture tested specimen.

FIG. 14 is a scanning electron micrograph taken at about 2500×magnification of a polished cross-section of Sample H near the rupturesurface.

FIG. 15 is a scanning electron micrograph taken at about 10,000×magnification of a polished cross-section of Sample H near the rupturesurface.

FIG. 16 shows total strain vs. time for a 1100° C. stress rupturespecimen at about 70 MPa tensile load in air.

FIG. 17a is an approximately 50× magnification optical photomicrographof a polished cross-section of a NICALON® fiber reinforced ceramicmatrix composite revealing the presence of several microcracks in thematrix material between adjacent fiber tows.

FIG. 17b is an approximately 50× magnification optical photomicrographof a polished cross-section of a fiber reinforced ceramic composite bodywhich shows how additions of silicon carbide particulate placed betweenadjacent plies of NICALON® fiber virtually eliminates these matrixmicrocracks.

FIGS. 18a and 18 b are isometric drawings of the graphite supportfixture of Example 15 loaded with fabric preforms and in the unloadedcondition, respectively.

FIG. 19 is an S-N plot showing the life of a fiber reinforced ceramiccomposite body as a function of temperature and the maximum appliedcyclical tensile stress.

FIG. 20 is a plot of sample strain versus time for a fiber reinforcedceramic composite body subjected to thermal cycling under an appliedtensile dead load.

FIG. 21 shows the four point flexural strength of a fiber reinforcedceramic composite body as a function of the atomic percentage (ratio) ofsilicon to boron in the reactant gases used to deposit by CVD a modifiedboron nitride coating layer on the fibrous reinforcement of thecomposite body.

DETAILED DESCRIPTION OF THE INVENTION AND PREFERRED EMBODIMENTS

In accordance with the present invention, there is disclosed a pluralityof distinct, but combinable, mechanisms for preventing undesirableoxidation (i.e., oxidation protection mechanisms) of one or morecomponents in composite bodies formed by various techniques.

By way of review, in composite material systems, particularly ceramicmatrix composite systems, frequently it is desirable for thereinforcement phase to debond and pull away or pull out of the matrix,at least partially. Such debonding and pull out absorbs mechanicalenergy which might otherwise have gone into fracturing the compositebody. Typically one or more coatings are applied to the reinforcementmaterial, e.g., the fibers, to accomplish the debonding under appliedload. Not only are the composite fabrication conditions (e.g., matrixdevelopment) harsh from a chemical corrositivity point of view, so arethe end use conditions, generally. Corrosion of the reinforcement or thedebond coating(s) becomes a concern because chemical reaction ordinarilyrenders the reinforcement or the debond coating(s) less effective. Thus,the concept of the duplex coating was developed: a debond coating on afiber itself coated with (and thereby protected by) a refractorymaterial.

Suitable ceramic filler materials which may be used in the inventioninclude metal oxides, borides, carbides, nitrides, silicides, andmixtures or combinations thereof, and may be relatively pure or containone or more impurities or additional phases, including composites ofthese materials. The metal oxides include, for example, alumina,magnesia, calcia, ceria, hafnia, lanthanum oxide, neodymium oxide,samaria, praseodymium oxide, thoria, urania, yttria, beryllium oxide,tungsten oxide and zirconia. In addition, a large number of binary,ternary, and higher order metallic compounds such as magnesium-aluminatespinel, silicon aluminum oxynitride, borosilicate glasses, and bariumtitanate are useful as refractory fillers. Additional ceramic fillermaterials may include, for example, silicon carbide, silica, boroncarbide, titanium carbide, zirconium carbide, boron nitride, siliconnitride, aluminum nitride, titanium nitride, zirconium nitride,zirconium boride, titanium diboride, aluminum dodecaboride, and suchmaterials as Si—C—O—N compounds, including composites of thesematerials. The ceramic filler may be in any of a number of forms, shapesor sizes depending largely on the matrix material, the geometry of thecomposite product, and the desired properties sought for the endproduct, and most typically are in the form of whiskers and fibers. Thefibers can be discontinuous (in chopped form as staple) or in the formof a single continuous filament or as continuous multifilament tows.They also can be in the form of two- or three-dimensional wovencontinuous fiber mats or structures. Further, the ceramic mass may behomogeneous or heterogeneous.

In a major aspect of the present invention, the oxidation protectionmechanisms of the invention include the use of getterer materials whichare present in at least a portion of the composite body (e.g., in atleast a portion of the matrix; in at least a portion of one or moreinterfacial coatings; or in, on or adjacent to at least a portion of thereinforcing materials, etc.). These getterer materials tend to scavenge(e.g., react with) undesirable oxidants which enter the composite bodysuch as, for example, by diffusion mechanisms, through microcracks, etc.These oxidation protection mechanisms may, in certain embodiments, alsoinclude techniques for reducing the number and/or size of suchmicrocracks in a portion of or throughout the composite body. Thereduction in the amount and/or size of microcracks may limit the abilityof undesirable oxidants to negatively impact the reinforcementmaterial(s) in the composite body.

When a composite body is put into service in an oxidizing environment,and assuming that the oxidizing environment would have an adverse effectupon the reinforcing material, some type of oxidation protectionmechanism should be utilized to prevent the reinforcement from oxidizingundesirably. If a getterer material was placed on, or at least in closeproximity to, the reinforcing material, then an oxidant which came intocontact with the getterer material such as, for example, by diffusionmechanisms through microcracks, etc., could be gettered (e.g., reacted)by the getterer materials, thereby ameliorating undesirable reaction(s)with the reinforcing material. Further, if the getterer material forms aparticular compound, such as for example, a glass, the compound couldprovide even further oxidation protection to the reinforcing material.In this regard, if a formed glass had an appropriate viscosity, then theformed glass could flow into any microcracks which may be present nearthe formed glass, thereby permitting the formed glass to function as acrack sealant. Such a desirable compound is sometimes referred to as a“trap sealant.” In this regard, the formed glass should have an oxidantpermeability which is low enough to provide for suitable oxidationprotection of the composite body at the intended operation temperaturesfor a desirable amount of time (e.g., the intended lifetime of thecomposite body).

In another embodiment of the invention, a glassy material,glass-network-forming material or glass modifier material is provided toa composite body during fabrication. For example, one can envisioncoating woven ceramic fiber plies with a particulate slurry comprising aglass-former such as silica and, optionally, one or more structuralmodifiers such as alumina, zirconia, calcia, etc.

A number of candidate getterer materials useful in combination withvarious matrices and reinforcements will become apparent to an artisanof ordinary skill upon review of this disclosure. Specifically, in apreferred embodiment of the invention, many reinforcement materials(e.g., fibers) are susceptible to oxidation by oxidants such as oxygen.Accordingly, it often is vitally important to prevent oxygen fromcontacting the reinforcing fibers so as to prevent any negative effectsupon the fibers. In this regard, oxygen typically is transported to afiber surface by a combination of different mechanisms. In general,oxygen usually enters the surface of a composite body due to some flawpresent on the surface (e.g., machining marks, a broken or cracked outerprotective skin, etc.). Once the oxygen has permeated the surface of acomposite body, oxygen may then ingress further into the composite bodyby various channels present in the composite body due to microcrackingfrom processing, thermal shock, physical shock, etc. In addition,molecular or atomic oxygen diffusion may also occur in combination withthe physical ingress of oxygen into the composite body. If anappropriate oxygen getterer material was positioned such that the oxygenwhich ingressed into the composite could be gettered (e.g., reactedwith) by the oxygen getterer, then further ingress of that particularoxygen molecule would be inhibited. However, if additional oxygeningressed into approximately the same area in the composite, at somepoint substantially all of the oxygen gettering material will eventuallyreact with the ingressing oxygen. At that point, it would be desirablefor another oxidation protection mechanism to occur. In this regard, ifthe oxygen gettering material were chosen so that one or more desirablecompounds (e.g., oxides or glasses) were formed upon a reaction with theoxygen, then such glasses or other oxides could block (e.g., flow into)any cracks, channels, microcracks, etc., to inhibit the physicaltransport of oxygen further into the composite body.

Examples of materials which function as suitable oxygen getterers andglass formers are boron, silicon, and the carbides and nitrides of boronand silicon. When reacted with oxygen the boron containing species mayform a boria based glass and the silicon containing species may form asilica based glass. Moreover, it is possible that when both boron oxideglass and silicon dioxide glass are present, the glasses may existindependently and/or may form a borosilicate glass. Still further, ifadditional materials are present in the vicinity of the forming glasses,such as aluminum (e.g., as a metal or an oxidized compound such asAl₂O₃) and/or zirconium, in various forms both oxidized andnon-oxidized, etc., it is possible to form in addition to those glassesmentioned above, glasses such as zirconium borosilicates, aluminumborosilicates, etc.

Thus, it should be apparent that one or more oxygen gettering materialscan be included in a composite body to form a number of desirablecompounds, such as those glasses discussed immediately above. In thisregard, it is possible to design a composite body so that when acomposite body is subjected to use in an oxidizing environment, a firstglass, such as a low melting boron oxide or borosilicate glass, willform and protect the reinforcing material of the composite at lowtemperatures. As the temperature of the composite body is increased, itis possible to form more refractory or higher softening point glasseswhich may result in oxidation protection at even higher servicetemperatures. For example, a high viscosity or high softening pointglass such as a zirconium borosilicate may extend the service life of acomposite body to heretofore believed to be impossible times at elevatedtemperatures. It also may be necessary to provide oxidation protectionat intermediate temperatures. In this regard, it may be desirable toform a glass such as an aluminum borosilicate which would bridge the gapin service temperature between, for example, the lower viscosity boronoxide glasses and the higher viscosity glasses such as zirconiumborosilicate. As is apparent from the above discussion, the number ofcombinations of oxygen gettering materials which can form desirableglasses, which may or may not react with other materials in thecomposite body, is quite large.

Further, it should be apparent to an artisan of ordinary skill thatdesirable glasses need not be formed entirely from the action of oxygengetterers. Instead, the desired glass or its components (glass-formers,modifiers, etc.) can be incorporated into the composite body duringcomposite fabrication. In service, glassy particulates may fuse to oneanother and flow into cracks. Glass formers and modifiers may alloyand/or react to form the desired glass. The compounds formed by “spent”oxygen getterers may also participate to modify these glasses originallyformed without oxygen getterer involvement.

Further, an important criterion in selecting materials which function togetter oxygen is the viscosity and oxygen permeability of the glassymaterial which is to be formed or modified due to reaction or alloyingwith an oxidized or “spent” gettering material. For example, in asilicon carbide fiber reinforced aluminum oxide material, an oxygengettering material which could be coated onto the fibers andsubsequently form a glass may need to be such that the glass so formedhas an oxygen permeability of about 1×10⁻⁹ g-O₂/cm² sec in order for thecomposite body to survive a few hours. However, if it is important forthe composite body to survive thousands of hours, the oxygenpermeability may need to be even lower; for example, about 1×10⁻¹²g-O₂/cm² sec may be necessary. By way of comparison, a microcrack mayexhibit an effective permeability of 1×10⁻⁶ g-O₂/cm² sec or more. It isof course apparent that oxygen permeability is a function of temperatureand an artisan of ordinary skill would need to determine the preciseservice temperature or temperatures that a composite body would beexposed to during service to determine the best combination of oxygengettering and glass forming materials to be used to extend the usefullife of the composite body.

Another factor to consider in designing an oxygen getterer system whichpossesses glass sealing characteristics is the effect of moisture,particularly at elevated temperatures. Specifically,boron oxide (e.g.,B₂O₃) glass dissolves in water according to the formula

B₂O₃+3H₂O←-→2H₃BO₃

At elevated temperature (e.g., 900° C.) H₃BO₃ is in the vapor phase.Thus, exposure of B₂O₃ glass to water vapor at such temperatures causesthe volatilization of the former. The reactivity/solubility ofborosilicate glasses with water is much less than that of straight boronoxide glass. Thus, all other things being equal, it may be better todesign a materials system to produce borosilicate glasses thanunmodified boria glass.

In general, oxygen gettering materials which form boron oxide orborosilicate glasses provide for relatively low temperature oxidationprotection (e.g., less than about 600° C.); however, oxygen getteringmaterials which form a calcium aluminosilicate glass may provideintermediate temperature oxidation protection (e.g., about 600° C.-1200°C.); oxygen getterers that form silicate glasses may provideintermediate to high temperature oxidation protection (e.g., about 600°C.-1800° C.); oxygen gettering materials which form a zirconium silicateglass or zircon structure may provide high temperature oxidationprotection (e.g., about 1200° C.-1800° C.); and oxygen getteringmaterials which form zirconia and silica glasses may provide for veryhigh temperature oxidation protection (e.g., about 1800° C.-2200° C.).

Accordingly, it is apparent that a composite body can be engineered sothat one or more getterer materials are included in the composite bodysuch that one or more desirable compounds (e.g., glasses) are formed.Each of the getterer materials could react with one or more oxidants atdifferent temperatures and form one or more desirable compounds (e.g.,one or more desirable glasses) which may provide for differing amountsof oxidation protection at different temperatures. Accordingly, acomposite body could be produced which contained a plurality ofdifferent oxidation protection mechanisms, wherein each oxidationprotection mechanism was included to provide for desirable oxidationprotection at different service temperatures of the composite body.

One exemplary manner of placing an oxidant getterer onto a reinforcingmaterial or at least in close proximity thereto would be to dip, paintor spray an appropriate material onto at least a portion of thereinforcing material prior to matrix formation or onto at least aportion of another material which is in contact with the reinforcingmaterial. Alternatively, or in conjunction with such coating by dipping,painting or spraying, chemical vapor deposition (CVD) or chemical vaporinfiltration (CVI) techniques could be utilized to obtain one or morecoatings on at least a portion of, or in a preferred embodiment,substantially all of, a reinforcing material. For example, a firstcoating comprising boron nitride could be deposited onto a reinforcingmaterial by CVI. One or more oxidant getterer materials might then beapplied to the boron nitride coating by dip coating, for example, thecoated reinforcing material into a solution comprising, for example, thenitrates or acetates of silicon, aluminum, zirconium and/or yttrium,which dip coated reinforcing material could then be heated in a nitrogenatmosphere, for example, to convert the nitrates to nitrides. It wouldbe desirable for such coatings to be capable of surviving any matrixformation steps as well as providing in-service oxidation protection. Ifnecessary, one or more additional coatings comprising, for example,silicon carbide could then be applied, for example, by CVI to protectthe underlying coatings and reinforcing material from chemicaldegradation during subsequent processing.

CVD or CVI is a particularly desirable means of placing one or moreoxidant getterers in proximity to a reinforcing material. Specifically,the precise location of the oxidant getterer in relation to thereinforcing material may be highly controlled. For example, if a firstcoating comprising boron nitride is to be applied, an oxidant getterercomprising aluminum or silicon could be applied as aluminum nitride orsilicon nitride, respectively, using CVI. Further, the oxidant getterercould be applied before, during and/or after the deposition of one ormore coatings to the reinforcing material to produce a coatedreinforcing material having an oxidant getterer underneath, mixed within(e.g., intermixed) and/or on top of (e.g., exterior to) the coatings.Moreover, it is possible to apply different oxidant getterer materialsat different locations relative to the reinforcing material. Forexample, one may choose to deposit, for example, an aluminum nitrideoxidant getterer beneath a first coating comprising boron nitride and azirconium nitride oxidant getterer on top of this first coating and/oron top of a second coating comprising silicon carbide. Still further,two or more oxidant getterer materials may be simultaneouslyco-deposited using CVD or CVI, such as, for example, simultaneousdepositions of oxidant getterers comprising aluminum and zirconium astheir respective nitrides. Depending upon conditions and choice ofcoating materials to be deposited, it is even possible to simultaneouslydeposit one or more of the oxidant getterer materials with the coatingsfor the filler material. For example, oxidant getterers (e.g., aluminum,silicon, yttrium, zirconium, etc.) may be co-deposited during depositionof the boron nitride debond coating onto the reinforcing fillermaterial. Finally, through careful control of the reactant gasconcentrations, a graded or tailored concentration of one or moreoxidant getterer materials can be achieved within a coating.

Without wishing to be bound by any particular theory or explanation, ithas been observed that boron nitride doped with silicon exhibitsincreased oxidation resistance, particularly where moisture is alsopresent. A convenient technique for producing such silicon doped boronnitride is by CVD, specifically by providing boron, nitrogen and siliconsources. Seemingly the silicon would chemically react with the nitrogensource to produce silicon nitride. Silicon nitride and boron nitrideco-deposition is not thermodynamically favorable at low temperatures, soto achieve a significant presence of silicon in the boron nitridedeposit, the co-deposition may need to be conducted at hightemperatures, for example at or above 1200° C. Because reaction ratestend to increase with increasing temperature, the precipitation of solidreaction product is rapid at such temperatures. The rapid depositionrates may not pose a problem for coating single filaments or fiber tows.However, it may be difficult or impossible to uniformly coat a fabric orstack of fabrics or a three-dimensionally woven fiber preform under suchconditions without the bulk of the deposit residing on the exterior ofthe preform and potentially sealing off the interior regions.

From a uniformity of coating deposition standpoint, slow deposition isbetter than rapid deposition. Low reaction temperatures are conducive toslower deposition rates. As stated above, low deposition temperaturesare not conducive thermodynamically to co-depositing silicon nitridewith boron nitride. Fortunately, it is still possible to co-deposit afew percent of silicon along with the balance of boron nitride at thelow temperatures, (e.g., about 700° C. to 800° C.) although it is notclear if silicon nitride is in solid solution with boron nitride or evenif the silicon is present as silicon nitride. More fortunate still hasbeen the discovery that even small amounts of silicon co-deposited withboron nitride can have a large beneficial effect on the resistance toenvironmental degradation of this silicon modified boron nitridecoating.

According to the present invention, modified boron nitride coatings havebeen produced that comprise from about 0.5 atom percent to about 3 atompercent silicon. Electron microscopy reveals that the modified boronnitride coating features a plurality of regions or “domains” each about5 to 20 nanometers in size. The individual domains exhibit a lamellarstructure, but the spatial orientation of the lamellae is essentiallyrandom from one domain to another.

It should be understood that the thickness of any coating which may beapplied to a reinforcing material influences a number of differentproperties, including the mechanical properties of a composite body, atboth ambient temperature and elevated temperatures, as well as theamount of oxidation protection afforded the reinforcing material. Ingeneral, the thickness of coatings on fibers in ceramic matrix compositebodies, where the ceramic matrix composite bodies are to be subjected toelevated temperature environments, should be from a few tenths of amicron thick to a few tens of microns in thickness and even morepreferably about 0.2 to about 20 microns in thickness. Specifically, ifa fibrous reinforcing material is chosen, the thickness of the coatingon the fiber should be sufficient to permit fiber pull-out to occur.Thicknesses greater than a few tens of microns may result in adversedegradation of mechanical properties (e.g., a coating which is too thickmay cause a failure mode to change from one which is predominantly fiberpull-out to a different failure mode which could have an overallweakening effect on the composite body), whereas thicknesses less than afew tenths of a micron may not provide for adequate oxidation protectionof the underlying fibers and/or not permit fiber/matrix debonding tooccur (e.g., if a thickness of coating was too thin, fibers may bebonded too strongly to the matrix thus inhibiting fiber pull-outmechanisms from occurring). Accordingly, numerous considerations need tobe taken into account when selecting the thickness of one or morecoatings to be placed upon a fiber reinforcement in a composite body.

In another aspect of the present invention, and particularly in regardto forming a ceramic matrix composite body by a directed metal oxidationof a parent metal, it has been discovered that a useful filler materialor strengthening component for the ceramic matrix composite body shouldbe provided with two or more coatings. The first or inner coating isapplied to the filler as a continuous film or layer, and preferablyforms a bond with the filler. The second and any subsequent coatings aresuperimposed over an underlying coating and become attached or bondedtherewith as additional layers or stratum. Each coating is applied as asubstantially continuous layer, and each is in substantially continuouscontact with the underlying coating or filler in the case of the firstcoating. The bond formed between adjacent surfaces may be weak ornegligible in that there may be little or no adhesion or connection, butin the preferred embodiment there is a measurable or appreciable bondingor union between surfaces.

In the embodiment of the invention in which multiple coatings are calledfor, two coatings applied to the filler material are normallysufficient. In such a system utilizing a duplex coating, the coatingsare selected to provide adequate mismatch in bonding strengths so as toallow for debonding and pull-out upon application of stress. Also, theduplex coating is selected to provide protection against degradation ofthe filler, and the outer coating is selected to exhibit wettability ofmolten parent metal and to protect the inner coating from degradation orcorrosion in high temperature, oxidizing environments under theconditions of the matrix formation process. Also, a system using twocoatings rather than three or more, may be somewhat more advantageousfrom an economic standpoint.

Thus, the coatings are selected so as to be compatible with the fillermaterial, and to the process conditions for the manufacture of thecomposites. Also, the coatings should complement each other in achievingthe desired characteristics or properties. In a ceramic composite systemhaving incorporated therein a filler with a duplex coating, for example,the first and outer coatings are selected to provide an adequatemismatch in interfacial shear strength so that one of the three zonaljunctions is weak relative to the remaining zonal junctions to providerelative movement between the inner coating and the filler, or betweencoatings, or between the outer coating and the adjacent ceramic matrix.In this manner, debonding and pull-out should occur, thereby improvingor enhancing the fracture toughness of the ceramic composite body.

Debonding and pull-out is especially beneficial for filler materialshaving a relatively high length to diameter ratio, such as fibers,typically at least about 2:1 and more particularly at least 3:1. Fillermaterial with a low length to diameter ratio such as particles orspheres, characteristically exhibits crack deflection toughening.

In applying the coatings to the filler material, the thickness of eachcoating and the cumulative thickness of all coatings can vary over awide range. This thickness can depend on such factors as the compositionof each coating and their interaction, the type and geometry of thefiller, and the process conditions and, for example, the parent metalused in the manufacture of the composite. Generally, the cumulativethickness for the coatings should be sufficient to completely cover theceramic filler material and protect it from, for example, oxidationdegradation, attack from molten metal, and other corrosive environmentswhich may be encountered in employment of the finished composite. In thepreferred embodiment, the inner coating is compatible with the fillermaterial so as not to degrade its integrity, and further the innercoating can be selected to allow for debonding and pull-out or shear.The coating system is selected to be compatible with the matrixmaterial, especially the precursor for the matrix, and further thecoating system is selected so as to be capable of withstanding theprocess conditions used in the manufacture of the composites. While theinner coating may afford adequate protection against degradation of thefiller or allow for shear between this first coating and the filler, asecond or outer coating is selected to be compatible with the processconditions employed in the manufacture of the ceramic composite body, inthat it should be substantially inert and not degrade, and furthershould exhibit wettability to molten parent metal when serving as aprecursor to the ceramic matrix. Also, if the first coating or fiber issusceptible to attack and degradation by the process environment duringcomposite manufacture or by attack of oxidants diffusing through thematrix during actual service, the second or outer coating is chosen toprotect the inner coating or fiber from exposure to processingconditions and/or end use conditions (e.g., the inner coating mayfunction as an oxygen getterer material alone or in combination withother components of the composite body such as other coatings or othermaterials in the composite body). Thus, the coating system protects thefibers from degradation, as does one coating superimposed on another,and concomitantly provides for compatibility for matrix formation anduse, and for relative movement to allow for shear. By reason of thiscoating system, structural degradation of the composite components ismitigated thereby prolonging the useful life and performance of thecomposite, and the fracture toughness of the composite is improved.

If the surface of a fibrous filler material is very irregular andexhibits nodules, barbs, fibrils, projections, or protuberances, thefiber can mechanically interlock or bond with the adjacent surfaceincluding the adjacent coating or adjacent fiber thereby impeding orpreventing debonding and pull-out, which can be deleterious to theproperties of the composite. It therefore is desirable to provide acoating system which is sufficiently thick to completely cover theirregularities in the fibers. Again, when large numbers of fibers orfilaments are being coated at the same time, the coating cannot be sothick as to isolate the fibers in the middle of a bundle from those nearthe exterior.

The thickness and properties of the coatings may vary depending on thedeposition process and the filler material. In a duplex coating system,the thickness for each coating, as measured from the center of a fillermaterial body out normal to the surface of the body, typically may rangefrom about 0.05 to about 25 microns, preferably to about 10 microns, butthe innermost coating can be as thin as a single monolayer in order toseparate the second coating from the filler particle. The cumulativethickness for a coating system may be to about 25 microns, and morepreferably 2-10 microns. Usually a coating system having a thicknesswithin this range can be applied to the filler by conventional or knownmeans and will provide the desired properties described above.

It has been found that a number of coating compositions can be employedin the coating system of this invention. These compositions include themetal oxides, nitrides, borides and carbides, alkaline metal salts,alkaline earth metal salts, carbon, silicon, and the like. The choice ofcoating compositions will depend on the filler material, thecompatibility of coatings to each other, and the process conditions forthe manufacture of the ceramic composite. For example, silicon carbidefibers are a popular choice as filler in composites intended for use atelevated temperatures. In order to provide for debonding and pull-out,the silicon carbide fibers may be coated with boron nitride whichprevents a relatively strong bond between the coated fiber and thesurrounding matrix. However, boron nitride may be degraded by theoxidation reaction conditions associated with a directed metal oxidationprocess. Further, boron nitride may not be wet by certain metals, suchas aluminum or silicon, under the conditions of the matrix formationprocess by infiltration (e.g., directed metal oxidation, meltinfiltration, etc.), and therefore as an outer coating would tend tointerfere with the matrix formation. However, an inner coatingexhibiting little or no wettability by the infiltrant metal underprocess conditions can be advantageous. For example, the coating systemmay have pores or flaws, but the contact angle of the molten infiltrantmetal with the inner coating may preclude transport of the metal throughany pores or flaws in the inner coating and thereby protect the fillerfrom attack by molten metal. The presence of an additional wettableouter coating on the filler would then avoid impedance to the matrixformation process. Therefore, a suitable outer coating such as siliconcarbide is applied to the boron nitride coating to achieve compatibilitywith the forming process and to protect the boron nitride fromdegradation, such as by oxidation. Silicon carbide is, for example, wetby doped aluminum and relatively oxidation-resistant in an airenvironment at 1000° C., whereas boron nitride is typically not wet byaluminum, and is oxidation-prone, at this temperature. Further, the bondbetween the two coatings is weak relative to the other bonds therebyfacilitating debonding and pull-out of the fibers during fracture. Otheruseful coating compositions include, for example, titanium carbide,silicon, calcium silicate, calcium sulfate, and carbon as the innercoating, and silicon, silica, alumina, zirconia, zirconium nitride,titanium nitride, aluminum nitride, and silicon nitride as an outercoating. Other suitable compositions for the first and outer coatingsmay be selected for use with the ceramic filler material provided thesecoatings complement each other as in the manner described above.

A typical cross-sectional representation of the coated ceramic fillermaterial is shown in FIG. 1 (discussed below in greater detail). In thistypical example, the ceramic filler material comprising silicon carbidebears a first inner coating of boron nitride and an additional outercoating of silicon carbide, thus a duplex coating. One or moreadditional outer coatings may be provided depending on need. Forexample, an additional outer coating of titanium carbide may be appliedto the coating of silicon carbide.

Moreover, it may be desirable to provide dual or multiple duplexcoatings such as boron nitride/silicon carbide/boron nitride/siliconcarbide. This multiple coating scheme may result in desirable internaloxidation protection mechanisms. Specifically, as discussed above, theinterface between boron nitride and silicon carbide may function as azonal debond junction, thus increasing the fracture toughness of amaterial, as well as providing for oxidation protection. As discussedabove, the precise composition and combination of coatings depends on anumber of factors including the processing or manufacturing environmentfor the composite body as well as the environment into which thecomposite body will be placed.

Non-oxide ceramic materials tend to decompose in the presence of oxygenat elevated temperatures. This problem can be particularly acute forhigh surface-to-volume geometries such as that of a fiber. Whether bychoice of design or by circumstances, many commercially availablenon-oxide ceramic fibers contain at least minor amounts of impuritymaterials. For example, in the case of a stabilized silicon carbidefiber such as NICALON® fiber, the impurity materials comprise oxygen andnitrogen. These impurities can stabilize or have the effect ofpotentially stabilizing the NICALON® fiber as manifested by preserving asubstantial fraction of such a fiber's ambient temperature strength upto elevated temperatures. Specifically, the oxygen and/or nitrogen atomsoccupy positions between microcrystalline silicon carbide grains. Theeffect of these impurities is to increase high temperature tensilestrength and reduce high temperature creep. Nevertheless, leftunprotected at elevated temperatures, NICALON® silicon carbide fibereventually loses a substantial portion of its original or as-fabricatedtensile strength, even when the exposure is conducted in a non-oxidizingenvironment such as, for example, in an argon atmosphere. Concurrentwith this strength loss, a mass loss from the fiber and, in particular,the oxygen and/or nitrogen impurities, has also been observed, such massbeing lost through, for example, volatilization. Without wishing to bebound by any particular theory or explanation, it seems that in at leastone instance this loss of the “stabilizing” impurities permits thefurther crystallization and growth of crystals within the fiber. Thefiber then consists of an assemblage of crystallites or grains of, forexample, silicon carbide, having distinct grain boundaries. If suchgrains continue to grow upon continued high temperature exposure of thefiber, the tensile strength of such a treated fiber may decrease withrespect to the same fiber in its original form. This strength loss canbe attributed, at least in part, to the development and growth of grainboundaries which may have or exhibit a strength limiting defect.

In certain circumstances, the presence of impurity atoms has the effectof stabilizing a crystal structure or stabilizing an amorphous structureunder conditions in which such a structure would not normally be stable.In other circumstances, this stabilizing effect manifests itself byameliorating the growth of the grains making up the reinforcing materialwhich can occur when a crystalline material is maintained at elevatedtemperatures, specifically, temperatures which are at a substantialfraction of the material's melting point. Such grain growth typicallyhas a deleterious effect on the strength of the fiber reinforcement andthus the component material itself because the size of strength-limitingflaws is often proportional to the grain size of a material. Thus, bystabilizing the as-fabricated small grain size of a reinforcing materialat elevated temperature, strength losses may be reduced or eliminated.Impurity materials (which are typically located at grain boundaries) mayserve to “pin” the grain boundaries of the grains making up thereinforcing material, thus stabilizing the reinforcing material againstgrain growth, and thus strength loss, at elevated temperatures. In suchcircumstances, it is therefore desirable to maintain the presence ofsuch impurity materials within a reinforcing fiber. Furthermore, becausesome of these impurities may tend to volatilize out of certainreinforcing fibers at elevated temperatures, it may be desirable to havein place around each reinforcing fiber a coating which may serve toprevent the stabilizing impurity materials from such volatilizing.

In the case of a stabilized silicon carbide fiber such as NICALON®fiber, a boron nitride coating has been observed to help maintain theoxygen and/or nitrogen stabilizing impurity atoms within the fiberduring elevated temperature exposure of the fiber. Without wishing to bebound by any particular theory or explanation, it has been hypothesizedthat the nitrogen component in the boron nitride coating effectivelyestablishes, at elevated temperature, a localized, fiber-externalnitrogen atmosphere or nitrogen partial pressure. This nitrogenatmosphere or nitrogen partial pressure represents a steep concentrationgradient of nitrogen across the fiber/coating interface. Thisconcentration gradient is biased against diffusion of the nitrogenimpurity out of the fiber and thereby tends to maintain the nitrogenstabilizing impurity within the fiber. Moreover, oxygen in the NICALON®fiber, if such a fiber is left unprotected, would likewise diffuse outof the fiber, similarly resulting in a degration of the fiber'sstrength. With an adjacent boron nitride coating, however, the diffusingoxygen contacts the boron nitride and reacts to form a very thin (e.g.,nanometers thick) coating of boron oxide at the fiber/boron nitrideinterface. This thin coating of boria thereby appears to inhibit furtherdiffusion of oxygen from the fiber to the external environment for thesame reason that the boron nitride coating suppresses nitrogen diffusionfrom the fiber.

Accordingly, it may be possible to protect similarly other non-oxideceramic reinforcing fibers from the effects of prolonged exposure atelevated temperatures by coating the fibers with coating materialshaving at least one element in common with either a constituent of thebasic fiber material or a necessary impurity material located within thebasic fiber material. Further, it may be possible through application ofthe above-described concepts to improve the high temperature stabilityof fibers based upon oxide systems. For example, a fiber reinforcingmaterial may comprise grains of aluminum oxide and small amounts of oneor more impurity substances located at aluminum oxide grain boundariesto stabilize the aluminum oxide grains against elevated temperaturegrain growth. If the stabilizing impurity material tends to volatilizeat these temperatures, a coating comprising at least one elementalcomponent of the impurity material which is in sufficient proximity tosuch a reinforcing material may help to maintain the stabilizingimpurity material within the aluminum oxide fiber.

The amount or thickness of coating material applied to the fillermaterial is also of importance, especially in composite materials whosematrices are formed by infiltration. Coatings which are too thick maytend to hinder infiltration of matrix material by sealing or isolatingregions of the permeable mass or preform to be infiltrated, particularlythose regions toward the center of the body. Conversely a debond coatingwhich is too thin may not provide sufficient debonding or pull-out ofthe filler material reinforcement from the matrix. Similarly, aprotective coating which is too thin may not be sufficiently protective.

It has been discovered for the case of coating reinforcement materialssuch as fibers and, in particular, fibers comprising silicon carbide andoccupying about 35-38 percent of the bulk volume of a preform, thethickness of the boron nitride coating which optimizes both the ambientand elevated temperature flexural strength of the resulting aluminumoxide matrix composite body is between about 0.2 micron and about 0.5micron and preferably averages about 0.3 micron. The flexural strengthof the composite has been observed to decrease for boron nitride coatingthicknesses below about 0.2 micron due to stressing of the compositebody beyond its yield point (i.e., proportional limit). In other words,as the composite body is stressed, an insufficient number of fibers pullout of the matrix to relieve the increasing elastic strain energy.Likewise, the flexural strength has been observed to decrease in somecomposite bodies where the boron nitride coating thickness exceeds about0.5 micron due to the onset of a new failure mode, specifically, forexample, that of interlaminar shear. Thus, with regard to mechanicalstrength, there does not appear to be any benefit to applying boronnitride coatings thicker than about 0.5 micron.

There also exists an optimal range for the silicon carbide overcoatthickness. Specifically, for the above-described system comprising boronnitride and silicon carbide coated onto fabric plies of woven continuousfibers comprising ceramic grade NICALON® silicon carbide and occupyingabout 35-38 percent of the bulk volume of a preform, the nominalthickness of the silicon carbide coating for optimizing the flexuralstrength of the alumina matrix composite formed by directed metaloxidation has been found to be in the range from about 2.0 to 2.3microns. More particularly, nominal silicon carbide coating thicknessesthinner than about 1.75 microns yielded composites with fracturestrengths which were significantly below the fracture strength ofcomposites having coatings of the thicknesses discussed above. Withoutwishing to be bound by any particular theory or explanation, this lossof strength may result from the relatively thin silicon carbide coatingsinadequately protecting the underlying boron nitride and/or siliconcarbide fiber reinforcement materials from chemical attack. Suchchemical attack of the reinforcement materials may occur during theformation of the matrix phase of the composite and/or during subsequentexposure of the formed composite to undesirable oxidant(s) at elevatedtemperatures. Likewise, nominal silicon carbide coating thicknessesgreater than about 2.3 microns have also yielded flexural strengthlosses. Again, without wishing to be bound by any particular theory orexplanation, silicon carbide coatings having nominal thicknesses whichare greater than about 2.3 microns appear to “can” or seal-off the spacewithin and/or between the fiber plies. This “canning” could then resultin the creation of closed porosity which may prevent subsequentinfiltration of oxidation reaction product into such closed porosityduring the directed metal oxidation process, thereby yielding weaklybonded fiber plies, and thus mechanically weakened, composite body.

It should be understood that the fiber coating thicknesses discussedherein have been calculated from the total weight gains which the fiberpreforms experience during the fiber coating process which in most casesherein refers to a chemical vapor infiltration (CVI) coating process.For boron nitride, the “actual” coating thickness as measured fromphotomicrographs of fiber cross-sections agree well with the calculatedvalues, as each fiber is coated with a relatively thin layer of boronnitride. For the thicker silicon carbide coating however, the two valuesdiverge. The subsequent coating of silicon carbide, however, isrelatively much thicker, and as the silicon carbide coatings build up inthickness, they may come into contact with one another, particularlywhere individual filaments are in close proximity, as is shown in FIG.14a, for example. This merging of individual silicon carbide coatingshas the effect of isolating some parts of the filaments from furthercoating deposition. Hence, the actual thickness of a silicon carbidecoating on those portions of a fiber where such a coating actuallyexists is somewhat larger than the “nominal” silicon carbide thicknesscalculated from weight gain values. Further, as a result of the natureof the CVI coating process, the silicon carbide coating thickness issomewhat greater for fibers in the outer preform plies than in the innerplies. Thus, for the above-described preform system comprising about35-38 volume percent of approximately 15 to 20 micron diameter fibers, anominal or calculated silicon carbide thickness of about 2.3 micronscorresponds to an actual coating thickness of about 4-6 microns near theexterior of the preform. Similarly, a nominal thickness of about 1.5microns corresponds to an actual coating thickness of about 2 micronsnear the preform exterior.

As mentioned immediately above, the nature of the CVD or CVI coatingprocess is to deposit a thicker-than-average layer on the exteriorregions of the fiber preform and a thinner-than-average layer in thosezones more towards the center of the preform. For example, depending onthe permeability or “openess” of the weave and the thickness of thepreform, an “average” SiC coating thickness of 2-2.3 microns, based uponweight gain values may correspond to actual depositions (as measured bymicroscopy) ranging from about 4-6 microns at the preform exterior toonly about 0.5-1.0 micron at the very center of the preform.

A technique has been found, however, to at least partially amelioatethis effect. Specifically, and as discussed elsewhere, when the preformis assembled as a stack of fabric plies, utilizing plies having a more“open” weave as the plies at the closest to the exterior of the preformprovides the CVI or CVD reactant gases greater access to the interiorregions of the preform. For example, a fiber preform may be assembledusing “eight harness satin weave” (8 HSW) and 12 HSW fabric plies.Because the 12 HSW plies exhibit a “tighter” weave than do the 8HSWplies, to make the coating thickness as uniform as possible through thepreform, the 8 HSW plies should be placed toward the exterior of thepreform and the 12 HSW plies should be placed at the center. An artisanof ordinary skill will appreciate the many possible combinations among12 HSW, 8 HSW and plain weave fabrics and even three-dimensionally wovenfiber preforms. For example, a 3-D woven fiber preform may be sandwichedin between at least one pair of plain weave fabrics having a lowervolumetric loading of reinforcement material. In general, and for allother factors being equal, the permeability or “openness” of wovenfabric increases in the following order: 12 HSW, 8HSW, 5 HSW, plainweave. The permeability of 3-D woven preforms depends upon too manyother variables to allow a generalization to be made; however, for thesame substance (e.g., NICALON® fiber), the permeability of any given 3-Dwoven preform may be estimated vis a vis that of a 2-D fabric weave bycomparing the bulk densities of the respective fiber forms, in otherwords, by comparing the respective total porosities.

In yet another aspect of the invention, a technique for reducing thenumber and/or size of microcracks formed during composite formationand/or formed during composite service or use, may be provided.Microcracks may be undesirable because such microcracks may permit readyaccess of undesirable oxidants to the reinforcement material(s) whichcan result in degradation of some properties of the composite body.Specifically, microcracking of a matrix material located betweenadjacent plies of fiber tows/bundles (e.g., silicon carbide fiber) canbe reduced or possibly even eliminated by introducing into the matrixone or more materials having a relatively low coefficient of thermalexpansion (e.g., lower than that of the matrix material) such as, forexample, silicon carbide particulate. Thus, to practice this embodimentof the invention, an appropriate material or combination of materialscould be inserted between one or more fiber tows or between fiber layersto form a preform from a combination of fibers and particulate. Afterformation of the preform, a ceramic matrix comprising, for example, anoxidation reaction product could be formed.

When plies or sheets of woven silicon carbide fiber tows are utilized inconjunction with an aluminum oxide matrix and more particularly wherethe warp yarns of adjacent plies are oriented at ninety degrees to oneanother, microcracks in the matrix may result. In this right angleorientation especially, there are (inevitable) regions between adjacentfiber plies substantially unoccupied by reinforcement fibers. During thedirected metal oxidation process, these regions as well as any voidspaces between individual fibers within a fiber tow and betweenindividual fiber tows within a fiber ply, fill in with ceramic oxidationreaction product. It has been observed that the ceramic matrix materialbetween adjacent fiber plies may be particularly susceptible tomicrocracking. FIG. 17A, for example, is an approximately 50×magnification optical photomicrograph of a polished cross-section ofsuch a fiber reinforced ceramic matrix composite which shows severalsuch cracks 300 within the ceramic matrix material 302 located betweenadjacent plies of woven tows of the reinforcement fibers 304.

Without wishing to be bound by any particular theory or explanation,microcracking of the matrix material, particularly, those portions ofthe matrix occupying the space between adjacent fiber plies, may resultfrom a difference or mismatch in the local thermal expansion coefficientof the composite material, specifically between the region of thecomposite within and between fiber plies, respectively. Accordingly, theobserved matrix microcracking may occur at some point during the coolingof the infiltrated preform from the process temperatures to ambienttemperature. For example, the NICALON® silicon carbide fibers employedin fabricating the fiber plies making up the preform have a thermalexpansion coefficient of about 4 ppm/° C., whereas the aluminum oxideand aluminum alloy making up a typical matrix material of a compositebody formed by directed metal oxidation have thermal expansioncoefficients of about 8 and 23, respectively.

Several concepts have been advanced in an effort to ameliorate theadverse consequences of these matrix microcracks. In particular, it hasbeen hypothesized that thermal expansion mismatch stresses might bereduced if the composite body were made more isotropic by, for example,reducing the difference in fiber orientation between adjacent plies fromninety degrees to thirty or forty-five degrees. Another idea has been toreduce the (largely unoccupied) space between adjacent fiber plies inthe preform by using thinner plies or by clamping the assemblage ofplies more tightly together during chemical vapor infiltration.

In contrast to these concepts directed to reducing the amount of spacebetween adjacent plies is the concept of filling this space with anothermaterial (e.g., a filler material) whose thermal expansion coefficientis selected such that the local thermal expansion coefficient of thecomposite material between adjacent fiber plies is closer (in value) tothat of the composite material within the fiber plies after such plieshave been embedded with the ceramic matrix. According to this reasoningand in view of the thermal expansion coefficient of the fiberreinforcement being lower than that of the matrix, an appropriate fillermaterial for the space between the plies could include a filler whichhas a thermal expansion coefficient lower than that of the matrix. Onesuch low thermal expansion coefficient filler material which has beenshown to be effective in this regard is silicon carbide. Not only doesadding silicon carbide particulate between the fabric plies reduce thethermal expansion coefficient of the otherwise unreinforced aluminamatrix material by virtue of the rule of mixtures, but additionally thelower expansion silicon carbide bodies act to constrain the contraction(upon cooling from the processing conditions) of the higher expansionalumina matrix Although the morphology of the added silicon carbidewhich has been successfully employed (as discussed later herein) was inthe form of particulates, other forms such as platelets, whiskers orchopped fibers could also be expected to work effectively.

Further, for many of the applications contemplated for the fiberreinforced composite materials of the present invention, high thermalconductivity of the composite is a desirable attribute. The presence offlaws in a material such as cracks tends to reduce the thermalconductivity of the material. Thus, a reduction in the number and/orsize of microcracks in a composite body also has the desirable affect ofincreasing thermal conductivity. Moreover, it is possible to furtherincrease the thermal conductivity of the body through selection ofmaterials having relatively high thermal conductivity and a low thermalexpansion coefficient. In this regard, the selection of silicon carbidefor placement as a filler material between plies of silicon carbidefiber tows is also an excellent choice because of the relatively highthermal conductivity of silicon carbide.

The filler material may be introduced between the fiber plies eitherbefore or after the CVI coating of the fiber plies. In a preferredembodiment, however, the filler materials which are introduced to thefiber plies are, at some point, CVI coated. In a particularly preferredembodiment, the filler materials are introduced between the fiber pliesprior to CVI coating by means of, for example, vibration or slurryinfiltration. Thus, the plies are assembled into a preform in thepresence of such filler materials and the resulting preform assembly isthen subsequently CVI coated such that both the fibers comprising thetows, as well as the filler material between adjacent plies of fibertows are simultaneously coated.

Because increasing the thermal conductivity of a ceramic body reducesits susceptibility to cracking due to thermal shock, one may choose toapproach the matrix microcracking problem by selecting filler materialsfor placement between the fiber plies, not necessarily based upon lowthermal expansion coefficient, but based upon high thermal conductivity.Accordingly, even filler materials having relative high thermalexpansion coefficients such as, for example, TiB2, may be gainfullyemployed in this way to reduce that matrix microcracking which is due tothermal shock, provided that such filler materials possess high thermalconductivity.

To further expand on this fiber reinforcement embodiment, one or moresubstantially non-reactive filler materials different from the fibrousfiller material may be added among the fibers making up a preform (e.g.,a 3-D woven fiber preform) or between the layers of fabric making up apreform. Provided that sufficient fiber loading remains to accomplishtheir purpose (e.g., composite toughening), a wide variety of otherfiller materials may be added to tailor a host of desired properties,for example, bulk density, thermal conductivity, wear resistance,ballistic performance, etc.

For example, once the required toughness has been met by providing acertain volumetric loading of de-bondable fibers, one or more otherproperties may be tailored through the addition of a different fillermaterial. For example, it may be desirable to improve the hardness ofthe composite material through addition of hard filler materials such assilicon carbide, boron carbide, the transition metal carbides, titaniumdiboride and/or boron carbide.

One convenient method for incorporating such additional filler materialsis by means of slurry infiltration or impregnation. Specifically, thedifferent filler material may be provided in whisker or particulate formand sipersed in water, an organic solvent or a preceramic polymer suchas CERASET™ SN inorganic polymer (Lanxide Performance Materials, Newark,Del.) to make a slurry. The slurry could then be painted, sprayed orpoured onto fabric plies of the fibrous reinforcement. Alternatively,the fabric plies or 3-D fibrous preforms could be dipped into theslurry. Optionally, pressure or vacuum could be administered to assistthe infiltration of the slurry into the void space between the fibers.The slurry infiltrated fibrous preform or fabric plies (once assembledinto the desired shape of the final self-supporting body) would then beheated to a modest temperature sufficient to remove volatiles or cureany polymeric components.

It should be understood that while this disclosure relates primarily tomatrices which are formed by directed metal oxidation, the conceptsdisclosed should be applicable to other matrix/fiber combinations.Accordingly, both reduction of microcracking and the increasing ofthermal conductivity can be enhanced in other systems as well.

The first and outer coatings, typically, are deposited onto the ceramicfiller material by conventional or known means such as chemical vapordeposition, plasma spraying, physical vapor deposition, platingtechniques, sputtering or sol-gel processing. Achievement of asubstantially uniform coating system according to these prior arttechniques is within the level of skill in this art. For example,chemical vapor deposition of a uniform coating of boron nitride onceramic filler materials can be achieved by using boron trifluoride andammonia at a temperature of about 1000-1500° C. and a reduced pressureof 1-100 torr; boron trichloride and ammonia at a temperature of600-1200° C. and reduced pressure of 0.1-100 torr; borazine at atemperature of 300-650° C. and a reduced pressure of 0.1-1 torr; ordiborane and ammonia at a temperature of 600-1250° C. and a reducedpressure of 0.1-1 torr. A coating of silicon carbide by chemical vapordeposition can be accomplished, for example, by usingmethyltrichlorosilane at a temperature of 800-1500° C. and a pressure of1-760 torr; dimethyldichlorosilane at a temperature of 600-1300° C. anda reduced pressure of 1-100 torr; and silicon tetrachloride and methaneat a temperature of 900-1400° C. and a reduced pressure of 1-100 torr.

It should be understood that various combinations of ceramic materialshaving one or more coatings may be produced depending on the specificproperties desired in the coated ceramic material and its ultimateapplication. A possible combination includes silicon carbide fiber witha first layer of titanium carbide and an additional outer layer ofsilicon nitride. Another coating system includes silicon carbide fiberwith a first coating of boron nitride and additional outer coatings ofsilicon carbide and alumina.

In the manufacture of ceramic matrix composites according to thedirected metal oxidation embodiment of the invention, the coatedmaterials may be provided in the form of a loose mass or may be laid upinto a porous preform of any desired configuration. The parent metal isplaced adjacent the preform. The parent metal is then heated in thepresence of an oxidant to above its melting point whereby the moltenmetal oxidizes to form and develop an oxidation reaction productembedding the coated ceramic material. During growth of the oxidationreaction product, the molten parent metal is transported through its ownotherwise impervious oxidation reaction product, thus exposing freemetal to the oxidizing atmosphere to yield additional reaction product.The result of this process is the progressive growth of aninterconnected ceramic oxidation reaction product which optionally maycontain nonoxidized parent metal.

A variety of ceramic matrices may be produced by the oxidation reactionof parent metals depending upon the choice of parent metal and oxidant.For example, ceramic matrices may include oxides, nitrides, borides, orcarbides of such parent or infiltrant metals as aluminum, silicon,titanium, tin, zirconium or hafnium. The ceramic matrix composites ofthe invention may comprise, by volume, 5 to 85% of the coated ceramicfiller materials and 95 to 15% of ceramic matrix. A useful compositecomprises an alumina matrix formed by the oxidation reaction of aluminumparent metal in air, or an aluminum nitride matrix by oxidation reaction(i.e., nitridation) of aluminum in nitrogen, and incorporating as areinforcing filler such materials as alumina, silicon carbide, siliconnitride, etc., bearing the coating system.

The choice of parent metal and oxidant will determine the composition ofthe polycrystalline matrix, as explained in the Commonly Owned Patentsand Patent Applications. Thus a filler bering the coating system mayhave admixed therewith a solid or liquid oxidant, such as boron, silica,or glasses (e.g., low melting glasses), or the oxidant may be gaseous,such as an oxygen-containing gas (e.g. air) or a nitrogen-containing gas(e.g. forming gas typically comprising, by volume, 96% nitrogen and 4%hydrogen).

Another useful composite material system is that of melt infiltration.Here, a silicon-based metal is melted and contacted to a permeable mass.The permeable mass comprises a material which can be wetted by moltensilicon, such as silicon carbide. Under wetting conditions, moltensilicon-containing metal can infiltrate such a permeable mass in apressureless manner. The infiltration typically is conducted under aninert atmosphere such as argon, or in a vacuum. The permeable massoptionally may include a carbon source, typically graphite, which mayreact with the infiltrating silicon to form silicon carbide in thematrix. Depending upon the amount of carbon source and the degree ofreaction, the in-situ formed silicon carbide may be interconnected ordiscrete, discontinuous bodies.

The oxidation protection mechanisms of the present invention can also beapplied to composite systems whose matrices may be formed by chemicalvapor infiltration (e.g., CVI SiC) or by repeated infiltration andpyrolysis of ceramic precursor polymers such as the polysilazanes.

It should be understood that while this disclosure relates primarily tomatrices which are formed by directed metal oxidation, the conceptsdisclosed should be applicable to other matrix processing systems suchas sintering, hot pressing or other infiltration techniques employed toproduce glass (e.g., “Black Glass” or “Comp Glass”), metal, polymer orother ceramic matrices.

The following examples illustrate certain aspects and advantages ofvarious embodiments of the invention.

EXAMPLE 1

Two fiber-reinforced alumina-matrix ceramic composite bodies werefabricated in accordance with the present invention. The fibers employedwere NICALON® ceramic grade silicon carbide as Si—C—O—N (from NipponCarbon Co., Ltd., Japan) measuring approximately 2 inches long andapproximately 10-20 lm in diameter. Each fiber was coated via chemicalvapor deposition with a duplex coating. The duplex coating comprised a0.2-0.5 lm thick first coating of boron nitride applied directly to thefiber, and a 1.5-2.0 lm thick second (outer) coating of silicon carbideapplied to the boron nitride coating.

The duplex coated fibers were gathered into bundles, each containing 500fibers tied with a single fiber tow. Two, 2 inch square by ½ inch thickbars of aluminum alloy designated 380.1 (from Belmont Metals, having anominally identified composition by weight of 8-8.5% Si, 2-3% Zn, and0.1% Mg as active dopants, and 3.5% Cu, as well as Fe, Mn, and Ni, butthe actual Mg content was sometimes higher as in the range of0.17-0.18%) were placed into a bed of Wollastonite (a mineral calciumsilicate, FP grade, from Nyco, Inc.) contained in a refractory cruciblesuch that a 2 inch square face of each bar was exposed to the atmosphereand substantially flush with the bed, while the remainder of each barwas submerged beneath the surface of the bed. A thin layer of silicasand was dispersed over the exposed surface of each bar to serve as anadditional dopant. Three of the above-described bundles of duplex-coatedfibers were placed on top of each of the two sand-layered metalsurfaces, and these set-ups were covered with Wollastonite.

The crucible with its contents was placed in a furnace which wassupplied with oxygen at a flow rate of 500 cc/min. The furnacetemperature was raised to 1000° C. at a rate of 200° C./hour, and heldat 1000° C. for 54 hours.

The crucible was then removed while the furnace temperature was at 1000°C., and allowed to cool to room temperature. The ceramic compositeproducts were recovered. Examination of the two ceramic compositeproducts showed that an alumina ceramic matrix, resulting from oxidationof aluminum, had infiltrated and embedded the fiber bundles.

Two specimens were machined from each of the two ceramic compositeproducts. FIGS. 1 and 2 are scanning electron micrographs at about 350×magnification and about 850× magnification, respectively, showing thisceramic matrix composite. Referring to the micrographs, there is shownthe alumina matrix 2 incorporating silicon carbide fibers 4 bearing afirst inner coating 6 of boron nitride and an outer coating 8 of siliconcarbide. One machined specimen from each composite product was testedfor flexural strength (Sintec strength testing machine, Model CITS2000/6, from Systems Integrated Technology Inc., Stoughton, Mass.) in 4point bend with a 12.67 mm upper span and a 28.55 mm lower span. Thevalues obtained were 448 and 279 MPa. The remaining specimen from eachproduct was tested for Chevron notch fracture toughness, and the valuesobtained were 19 and 17 MPa-m^(½), respectively. FIG. 3 is a scanningelectron micrograph at 250× magnification of the fractured surface ofthe ceramic composite showing extensive pull-out of the fibers.

This run was repeated with the exception that the NICALON® fibers werenot coated. FIG. 4 is a scanning electron micrograph at 800×magnification of the fractured surface showing essentially no pull-outof the fibers. Typical values for strength ranged from 100-230 MPa, andfor toughness ranged from 5-6 MPa-m^(½.)

The utility of coated filler material made according to the invention isclearly demonstrated by this Example and the comparative data.

EXAMPLE 2

The following Example demonstrates a method for forming a fiberreinforced ceramic composite body, and illustrates the resultantmechanical properties of the body from about room temperature to about1400° C. Specifically, this Example demonstrates a method for forming asilicon carbide fiber reinforced alumina composite body wherein thesilicon carbide fibers are coated with a first layer of boron nitrideand a second layer of silicon carbide to create a debond zone betweenthe silicon carbide fiber and the alumina matrix.

A fabric preform 103 was made by stacking a plurality of layers of 8harness satin weave (8 HSW) fabric and 12 harness satin weave (12 HSW)fabric made from ceramic grade NICALON® silicon carbide fiber (obtainedfrom Dow Corning Corporation, Midland, Mich.) on top of each other.FIGS. 5a and 5 b are schematics depicting a top view and across-sectional view respectively of the as-is position for a HSWfabric. In reference to FIGS. 5a and 5 b, a HSW fabric is designated tobe in the “as-is position” when, as viewed in cross-section, the axes ofthe warp yarns 92 of the fabric 90 are in the plane of thecross-sectional view and are located at the bottom (i.e., as shown inthe cross-sectional view) of the fabric 90 and the axes of the fillyarns 91 are perpendicular to the plane of the cross-sectional view andare located at the top of the fabric 90. The orientation of additionalfabric layers can be described in reference to the as-is position. Forexample, as depicted in FIG. 5c, additional fabric layers can be (1)rotated about an axis 93 perpendicular to the plane of the fabric 90and/or (2) rotated about an axis 94 perpendicular to the plane of thecross-section of the fabric 90 and then subsequently contacted orlayered upon a fiber layer positioned in the as-is configuration. Thus,for example, as schematically depicted in cross-section in FIG. 5d, asubstantially square fabric preform 103 can be made from 8 pieces of HSWfabric, stacked in the following sequence:

A first fabric layer 95 comprising an 8 HSW fabric was placed on asupporting surface in the as-is position to start the fabric preform103;

A second fabric layer 96 comprising a 12 HSW fabric, was rotated about90° in the counterclockwise direction from the as-is position about anaxis 93 perpendicular to the plane of the fabric and was placed on thefirst fabric layer 95 so that the edges of the second fabric layer 96were substantially aligned with the edges of the first fabric layer 95;

A third fabric layer 97 comprising a 12 HSW fabric, in the as-isposition, was placed on the second fabric layer 96 so the edges of thethird fabric layer 97 were substantially aligned with the edges of thesecond fabric layer 96;

A fourth fabric layer 98 comprising a 12 HSW fabric, was rotated about90° in the counterclockwise direction from the as-is position about anaxis 93 perpendicular to the plane of the fabric and was placed on thethird fabric layer 97 so that the edges of the fourth fabric layer 98were substantially aligned with the edges of the third fabric layer 97;

A fifth fabric layer 99 comprising a 12 HSW fabric, was rotated about90° in the counterclockwise direction from the as-is position about anaxis 93 perpendicular to the plane of the fabric and then rotated about180° in the clockwise direction about an axis 94 perpendicular to theplane of the cross-sectional view of the fabric and was placed on thefourth fabric layer 98 so that the edges of the fifth fabric layer 99substantially aligned with the edges of the fourth fabric layer 98;

A sixth fabric layer 100 comprising a 12 HSW fabric, was rotated about180° in the clockwise direction from the as-is position about an axis 94perpendicular to the plane of the cross-sectional view of the fabric andwas placed on the fifth fabric layer 99 so that the edges of the sixthfabric layer 100 were substantially aligned with the edges of the fifthfabric layer 99;

A seventh fabric layer 101 comprising a 12 HSW fabric, was rotated about90° in the counterclockwise direction from the as-is position about anaxis 93 perpendicular to the plane of the fabric and then rotated about180° in the clockwise direction about an axis 94 perpendicular to theplane of the cross-sectional view of the fabric and was placed on thesixth fabric layer 100 so that the edges of the seventh fabric layer 101were substantially aligned with the edges of the sixth fabric layer 100;and

Finally, an eighth fabric layer 102 comprising an 8 HSW fabric, wasrotated about 180° in the clockwise direction from the as-is positionabout an axis perpendicular 94 to the plane of the cross-sectional viewof the fabric and was placed on the seventh fabric layer 101 so that theedges of the eighth fabric layer 102 were substantially aligned with theedges of the seventh fabric layer.

In reference to FIG. 5e, the fabric preform 103 comprising two 8 HSWouter fabric layers and six 12 HSW inner fabric layers and measuringabout 6.75 inch (171 mm) square and about 0.125 inch (3.2 mm) thick wasplaced on a perforated graphite plate 104 machined from Grade AXF-5Qgraphite (Poco Graphite, Inc., Decatur, Tex.) which measured about 7.75inches (197 mm) square and about 0.5 inch (13 mm) thick. The innerperforated region 105 of the perforated plate measured about 6.25 inches(159 mm) square. The holes 106 of the perforated region 105 had adiameter of about 0.25 inch (6.4 mm) and a center-to-center spacing ofabout 0.375 inch (9.5 mm) and comprised a 17 hole×17 hole array whichwas bordered by an about 1 inch (25 mm) unperforated region. After thefabric preform 103 had been placed on the first graphite plate 104, asecond graphite plate 104, substantially the same as the first, wasplaced over the fabric preform 103 and the plates were clamped usingC-clamps to compress the fabric preform 103. Two graphite channelmembers 107 machined from Grade AXF-5Q graphite (Poco Graphite, Inc.,Decatur, Tex.) and measuring about 7.75 inches (197 mm) long were placedover common ends of both perforated graphite plates 104 so as to contactopposite ends of the first and second perforated graphite plates 104thereby creating a preform containment fixture 108. FIG. 5e is anisometric schematic view of the preform containment fixture 108. Afterthe graphite channels 107 were secured to the perforated plates 104, theC-clamps were removed from the perforated plates 104 and the elasticforce exerted by the compressed fabric preform 103 biased the perforatedgraphite plates 104 against the graphite channel members 107 to form arelatively rigid preform containment fixture 108. The warp yarns 92 ofthe eighth layer 102 of the fabric preform 103 within the graphitecontainment fixture 108 were positioned so as to be parallel to thelength of the graphite channel members 107 of the preform containmentfixture 108.

The graphite containment fixture 108 containing the fabric preform 103was placed into a reactor chamber of a chemical vapor infiltrationapparatus having an outer diameter of about 12 inches (305 mm). Theinner diameter of the reactor chamber measured about 9.45 inches (240mm) after being lined with a quartz tube having a wall thickness ofabout 0.5 inch (13 mm) and lined with a graphite tube having a wallthickness of about 0.25 inch (6.4 mm). The warp yarns 92 of the eighthlayer 102 of the fabric preform 103 were parallel to the gas flowdirection within the chamber as well as being parallel to thelongitudinal axis of the reactor chamber. The reactor chamber was closedand evacuated to about 0.004 inch (0.1 mm) of mercury (Hg). Then thereactor chamber was heated to about 800° C. at about 10° C. per minuteso that the contents of the reactor chamber were at about 730° C., asindicated by a thermocouple contained therein. When the temperaturewithin the reactor chamber reached about 730° C., a gas mixturecomprised of ammonia (NH₃) flowing at about 1200 standard cubiccentimeters per minute (sccm) and boron chloride (BCl₃) flowing at about800 sccm was introduced into the reactor chamber while maintaining atotal operating pressure of from about 0.047 to about 0.051 inches ofmercury (about 1.2 to about 1.3 mm Hg). After about 6.5 hours at about730° C., the gas mixture flowing into the reactor chamber wasinterrupted, the power to the furnace heating the reactor chamber wasinterrupted, and the furnace and its contents were naturally cooled toabout 200° C. At about 200° C., the reactor chamber door was opened andthe graphite containment fixture 108 was removed, cooled anddisassembled to reveal that the fibers of the fabric layers of thefabric preform 103 were coated and that the fabric layers comprising thefabric preform 103 were bonded together by a boron nitride coatingformed during the process at about 730° C., thereby forming a coated andbonded fabric preform 109. The boron nitride coating had a thickness ofabout 0.4 microns.

The boron nitride coated and bonded fabric preform 109 was thensuspended from a graphite cantilever support fixture 110 made from GradeAXF-5Q graphite (Poco Graphite, Inc., Decatur, Tex.) by wires 111comprised of a Kanthal® iron-chromium-aluminum alloy all of which aredepicted schematically in FIG. 5f. The graphite cantilever supportfixture 110 and the boron nitride bonded fabric preform 109 were thenreplaced into the reactor chamber of the chemical vapor infiltrationapparatus discussed above such that the warp yarns 92 of the eighthlayer 102 comprised of the 8 harness satin weave fabric were parallel tothe gas flow direction within the chamber as well as being parallel tothe longitudinal axis of the reactor chamber. After the reactor chamberdoor was closed, the reactor chamber and its contents were evacuated toabout 0.591 inches (15 mm Hg) and hydrogen gas flowing at about 2500sccm was introduced into the reactor chamber. The reactor chamber washeated at about 10° C. per minute so that the contents of the reactorchamber were at about 925° C. as indicated by a thermocouple therein.When the reactor chamber contents were at about 925° C., additionalhydrogen, flowing at about 2500 sccm, was introduced into the reactorchamber to give a total hydrogen gas flow rate of about 5000 sccm. Oncethe temperature of the contents of the reactor chamber had substantiallycompletely stabilized at about 925° C., about 2500 sccm hydrogen werediverted away from direct entry into the reactor chamber, and were firstbubbled through a bath of trichloromethylsilane (CH₃SiCl₃) also known asmethyltrichlorolsilane (MTS) (Hulls/Petrarch System, Bristol, Pa.),maintained at about 25° C., before entering the reactor chamber. Afterabout 26 hours at about 925° C., the power to the furnace heating thereactor chamber was interrupted and the about 2500 sccm hydrogen thatwas being directed through the MTS bath was again permitted to flowdirectly into the reactor chamber to reestablish a direct hydrogen gasflow rate of about 5000 sccm into the reactor chamber. It was noted thatabout 4.75 liters of MTS had been consumed during the 26 hour of the runat about 925° C. After about a half hour during which a hydrogen gasflow rate at about 5000 sccm was maintained, the hydrogen flow rate wasinterrupted and the furnace and its contents were evacuated to about0.039 inches 0.1 mm of mercury (Hg). The pressure within the reactorchamber was then allowed to increase to about atmospheric pressure whileargon was introduced at a flow rate of about 14 liters per minute. Afterthe reaction chamber had cooled to a temperature of about 200° C., theargon flow rate was interrupted and the reaction chamber door wasopened. The graphite cantilever support fixture 110 and the fabricpreform were removed from the reactor chamber to reveal that the boronnitride bonded fabric preform 109 had been coated with a second layer ofsilicon carbide thereby forming a silicon carbide (SiC)/boron nitride(BN)-coated fabric preform 112. The silicon carbide had an overallaverage thickness of about 2.3 microns, as calculated from the weightgain of the preform during the silicon carbide coating procedure, asalluded to previously.

A wax box pattern having a closed end and outer dimensions of about 7inches (178 mm) square by about 2 inches (51 mm) tall and a wallthickness of about 0.25 inches (6.5 mm) was assembled from hightemperature wax sheet (Kit Collins Company, Cleveland, Ohio) whichcontained adhesive backing on one side thereof. The wax box pattern wasassembled by using a hot wax knife. The closed end of the wax patternwas beveled at an angle of about 22°. A slurry mixture comprised byweight of about 5 parts BLUONIC® A colloidal alumina (BuntrockIndustries, Lively, Va.) and about 2 parts −325 mesh (average particlediameter less than about 45 μm) wollastonite (a calcium silicatemineral) was made by hand mixing the materials together. The slurrymixture was then painted onto the outer surface of the wax box patternwith a one inch sponge brush and wollastonite powder (−10, +100 mesh)having substantially all particles between about 150 and 2000 microns indiameter was sprinkled liberally onto the slurry mixture coating toprevent runoff and to form a first precursor layer of a shell 120. Thisprocedure was repeated to build additional layers of coating with anabout 0.5 hour drying period between the formation of the precursorlayers. When enough precursor layers of slurry mixture/coarsewollastonite were formed to produce a thickness of about 0.25 inch (6.4mm), the coated wax box pattern was set aside to dry at about roomtemperature for about 24 hours. The about 0.25 inch (6.4 mm) thickcoating nominally comprised about 12 slurry mixture/coarse wollastonitelayers. After the coated wax box pattern had substantially completelydried at about room temperature, the wax box pattern was placed into anair atmosphere furnace maintained under an exhaust hood and the furnaceand its contents were held at a temperature of about 120° C. for about 6hours, during which time the wax melted leaving behind an unfiredprecursor to an alumina bonded wollastonite shell 120. The furnace andits contents were then heated to about 950° C. in about 2 hours and heldat about 950° for about 4 hours to substantially completely remove anyresidual wax and ensure the sintering of the alumina bonded wollastoniteshell. The furnace and its contents were then cooled to about roomtemperature.

About 40 grams of VASELINE® petroleum jelly vehicle (Cheseborough Ponds,Inc., Greenwich, Conn.) were melted in a small aluminum weighing dish ona hot plate set at about medium heat until the jelly turned to a liquid.A clean sable brush was then used to substantially completely coat oneof the 6.75 inch (171 mm) square surfaces of the SiC/BN-coated fabricpreform 112 to provide an interface for the application of a nickeloxide powder. A mixture comprising about 8 grams of −325 mesh (particlediameter less than about 45 μm) nickel oxide powder and about 16 gramsof ethanol was applied with a sponge brush to substantially completelycover the petroleum jelly coated surface of the SiC/BN-coated fabricpreform. After the ethanol had substantially completely evaporated, theSiC/BN-coated fabric preform 112 was inserted into the alumina bondedwollastonite shell 120 such that the uncoated side of the SiC/BN-coatedpreform 112 not coated with the nickel oxide powder contacted the bottomof the shell 120, as shown in FIG. 5g. The spaces between the perimeterof the SiC/BN-coated fabric preform 112 and the walls of the aluminabonded wollastonite shell 120 were filled with coarse (−10, +100 mesh)wollastonite until the surface of the wollastonite powder wassubstantially flush with the nickel oxide powder-coated surface of theSiC/BN-coated fabric preform 112. The alumina bonded wollastonite shell120 containing the SiC/BN-coated fabric preform 112 was then placed ontostilts 122, which were made from fire brick, and was thereaftersurrounded by wollastonite powder 123 which was contained in arefractory boat 124. The SiC/BN-coated fabric preform 112 was thenleveled. About 1600 grams of a parent metal was distributed into four 30gram clay crucibles (obtained from J.H. Berge, Inc., South Plainfield,N.J.) in amounts of about 400 grams per crucible. The parent metalcomprised by weight of about 8.5 to 11.0 percent silicon, 3.0 to 4.0percent copper, 2.7 to 3.5 percent zinc, 0.2 to 0.3 percent magnesium,≦0.01 percent calcium, ≦0.10 percent titanium, 0.7 to 1.0 percent iron,≦0.5 percent nickel, ≦0.5 percent manganese, ≦0.35 percent tin, ≦0.001percent beryllium, ≦0.15 percent lead and the balance aluminum. Therefractory boat 124 and its contents, as well as the four 30 gram claycrucibles containing the parent metal, were placed into an airatmosphere furnace and the furnace door was closed. The furnace and itscontents were then heated from about room temperature to about 700° C.at about 400° C. per hour, during which time the VASELINE® petroleumjelly volatilized and the nickel oxide powder 125 fell onto the surfaceof the SiC/BN-coated fabric preform 112. After about an hour at about700°, during which time the parent metal 126 had substantiallycompletely melted, the parent metal 126 was then poured into the aluminabonded wollastonite shell 120 and onto the nickel oxide powder-coatedside of the SiC/BN-coated fabric preform 112, thereby covering thesurface of the preform 112. Wollastonite powder 127 was then poured ontothe surface of the molten parent metal 126 within the alumina bondedwollastonite shell 120 to substantially completely cover the surface ofthe molten parent metal. This assembly formed the lay-up for growth of aceramic matrix composite body. The furnace and its contents comprisingthe lay-up were then heated to about 950° C. in about an hour. Afterabout 90 hours at about 950° C., the furnace and its contents werecooled to about 700° C. in about 2 hours. At about 700° C., the lay-upwas removed from the furnace and residual molten parent metal wasdecanted from the alumina bonded wollastonite shell 120, the shell 120was quickly broken away from the SiC/BN-coated fabric preform 112 andthe preform 112 was buried in a silica sand bed to cool the preform 112to about room temperature. At about room temperature, it was observedthat an oxidation reaction product had grown into and substantiallycompletely embedded the SiC/BN-coated fabric preform 112, therebyforming a fiber reinforced ceramic composite body 130 having a pluralityof fabric layers comprised of harness satin weaves. Specifically, thefiber reinforced ceramic composite body 130 comprised two outer layersof 8 harness satin weave silicon carbide fabric and six inner layers of12 harness satin weave silicon carbide fabric embedded by an aluminumoxide oxidation product. The composite body also comprised a metallicconstituent comprising residual unreacted parent metal.

Once the ceramic composite body had been manufactured, a metal removalprocess was begun for the purpose of removing this residual parent metalwithin the composite body. The first step of the metal removal processwas to form a filler material mixture for infiltration by metalcontained in the formed ceramic matrix composite body.

Specifically, filler material mixture comprising by weight of about 90percent E67 1000 grit (average particle diameter of about 5 μm) alumina(Norton Co., Worcester, Mass.) and about 10 percent −325 mesh (particlediameter less than about 45 μm) magnesium powder (Reade ManufacturingCompany, Lakehurst, N.J.) was prepared in a one gallon NALGENE® widemouth plastic container (Nalge Co., Rochester, N.Y.). Alumina millingballs were added to the filler material mixture in the plastic containerand the container lid was closed. The plastic container and its contentswere placed on a jar mill for about 4 hours to mix the alumina andmagnesium powders together. After the alumina mixing balls had beenseparated from the alumina-magnesium filler material mixture 131, thefiller material mixture 131 was complete.

A stainless steel boat 132 measuring about 7 inches (179 mm) square byabout 2 inches (50.8 mm) deep and having a wall thickness of about 0.063inches (1.6 mm) was lined with a graphite foil box 133 made from a pieceof GRAFOIL\ graphite foil (Union Carbide Corp., Carbon ProductsDivision, Cleveland, Ohio). About 1 inch (25 mm) of the filler materialmixture 131 was hand packed into the bottom of the graphite foil linedstainless steel boat 132. The fiber reinforced ceramic composite body130 was then placed onto and forced into the filler material mixture131. Additional filler material mixture 131 was then poured over thefiber reinforced ceramic composite body 130 to substantially completelycover it. The filler material mixture 131 was then hand packed to ensuregood contact between the filler material mixture 131 and the fiberreinforced ceramic composite body 130, thereby forming a metal removallay-up as depicted schematically in cross-section in FIG. 5h.

The metal removal lay-up comprising the stainless steel boat 132 and itscontents was then placed into a resistance heated controlled atmospherefurnace and the furnace chamber door was closed. The furnace chamber andits contents were first evacuated to at least 30 inches (762 mm) ofmercury (Hg) vacuum, then the vacuum pump was disconnected from thefurnace chamber and nitrogen was introduced into the chamber toestablish about atmospheric pressure of nitrogen in the chamber. Thisoperation was repeated. After the pressure in the furnace chamberreached about atmospheric pressure, the furnace chamber and its contentswere heated from about room temperature to about 750° C. at a rate ofabout 250° C. per hour and held at about 750° C. for about 5 hours andcooled from about 750° C. to about 300° C. at about 200° C. per hourwith a nitrogen gas flow rate of about 4000 sccm being maintainedthroughout the heating and cooling. At about 300° C., the nitrogen flowwas interrupted, the furnace door was opened, and the stainless steelboat and its contents were removed and cooled by forced convection. Atabout room temperature, the filler material 131 was separated from thefiber reinforced ceramic composite body 130 and it was noted that themetallic constituent of the fiber reinforced ceramic composite body 130had been substantially completely removed. The fiber reinforced ceramiccomposite body 130 was then subjected to grit blasting by a sand blasterwhich operated with a working pressure of about 75 pounds per squareinch to remove any excess filler material that had adhered to thesurface of the composite body 130. The fiber reinforced ceramiccomposite body was then cut with a diamond saw and machined intomechanical test specimens measuring about 2.4 inches (60 mm) long byabout 0.2 inch (6 mm) wide by about 0.11 inch (3 mm) thick formechanical properties measurements, specifically flexeral strengthtesting.

Several of the machined mechanical test specimens were then subjected toadditional heat treatments. Except as otherwise noted, these heattreatments were limited to the fiber reinforced ceramic compositematerial of the present Example. Specifically, a first group of sampleswas heat treated at about 1200° C. for about 24 hours and a second groupof samples was heated treated at about 1200° C. for about 100 hours. Theheat treatments were effected by placing the mechanical test specimensonto alumina trays with the tensile side of the test specimen facingaway from the alumina trays. The alumina trays and their contents werethen placed into air atmosphere furnaces and heated to about 1200° C. ata rate of about 200° C. per hour. After about 24 hours at about 1200°C., the furnace containing the first group of samples was cooled toabout room temperature at a rate of about 200° C. per hour, whereasafter about 100 hours at about 1200° C., the furnace containing a secondgroup of samples, was cooled to about room temperature at a rate ofabout 200° C. per hour.

The flexural strengths of the fiber reinforced ceramic composite testspecimens were measured using the procedure defmed by the Department ofthe Army's proposed MIL-STD-1942A (Nov. 21, 1983). This test wasspecifically designed for strength measurements of high-performanceceramic materials. The flexural strength is defmed in this standard asthe maximum outer fiber stress at the time of failure. Afour-point-¼-point flexural test was used. The height and width of thetest bars were measured with a precision of about 390 microinch (0.01mm). The test bars were subjected to a stress which was applied at fourpoints by two lower span bearing points and two upper span bearingpoints. The lower span bearing points were about 1.6 inches (40 mm)apart, and the upper span bearing points were about 0.79 inch (20 mm)apart. The upper span was centered over the lower span, so that the loadwas applied substantially symmetrically on the test bar. The flexuralstrength measurements were made with a Sintec Model CITS-2000/6universal testing machine (Systems Integrated Technology, Inc.,Stoughton, Mass.). The crosshead speed during testing was about 0.02inch per minute (0.55° C., about 1300° C. and about 1400° C. wereperformed with another universal testing machine equipped with an airatmosphere resistance heated furnace (Advanced Test Systems, Butler,Pa.).

Table I contains a summary of the four point flexural strengths forNICALON® silicon carbide reinforced alumina oxidation reaction productcomposite bodies. Specifically, Table I summarizes the sample condition,the test temperature, the number of samples tested, the average flexuralstrength and standard deviation, the maximum flexural strength and theminimum flexural strength. These data suggest that the flexural strengthof fiber reinforced ceramic composite bodies subjected to the methods ofthe instant invention are substantially unaffected by test temperaturebetween about room temperature and about 1200° C. Moreover, these datasuggest that the flexural strengths of fiber reinforced ceramiccomposite bodies subjected to the methods of the instant invention areonly slightly degraded at test temperatures greater than 1200° C. and byextended exposure times at 1200° C.

EXAMPLE 3

This Example illustrates that fiber reinforced ceramic composite bodieshaving varying ceramic matrix composition can be formed. Specifically,Sample A of this Example comprised a silicon carbide fiber reinforcedalumina composite body; and Sample B of this Example comprised a siliconcarbide fiber reinforced aluminum nitride composite body.

Sample A

A SiC/BN-coated fabric preform measuring about 3.0 inches (76 mm) longby about 3.0 inches (76 mm) wide by about 0.125 inch (3.2 mm) thick wasprepared by stacking eight layers of 12-harness satin weave (12 HSW)fabric comprising silicon carbide fibers (ceramic grade NICALON® fibersobtained from Dow Corning Corporation, Midland, Mich.) the fibers havinga diameter ranging from about 394 microinch (10 μm) to about 787microinch (20 μm). The 12 HSW silicon carbide fabrics were stacked suchthat each succeeding fabric layer was placed with its fill yarns beingrotated about 90° with respect to the fill yarns of the previous fabriclayer. The fabric preform comprising the stacked layers were then placedinto a chemical-vapor-infiltration (CVI) reactor and the fibers werecoated with a first layer of boron nitride (BN) substantially inaccordance with the methods of Example 2. Thereafter, the reactionconditions in the CVI reactor were modified such that a CVI coating ofsilicon carbide (SiC) was placed on top of the BN coating substantiallyin accordance with the method of Example 2. The CVI coatings held thestacked fabric layers together, thereby forming the SiC/BN-coated fabricpreform.

The SiC/BN-coated fabric preform comprising the eight stacked layers of12 HSW fabric coated with a first layer of BN and a second layer of SiCwas placed into the bottom of a porous castable refractory boat havingholes at the bottom to facilitate air flow to the composite duringcomposite growth, thereby forming a lay-up. Specifically, the porouscastable refractory boat having an inner cavity measuring about 3.25inches (83 mm) square by about 3.0 inches (76 mm) deep and having a wallthickness of about 0.125 inch (3.2 mm) was cast from a mixture comprisedby weight of about 56.3% plaster of Paris (BONDEX®, BondexInternational), about 28.1% water and about 15.6% 90 grit alumina (E1ALUNDUM®, Norton Company, Worcester, Mass.). After the SiC/BN-coatedfabric preform was placed into the porous castable refractory boat, −325mesh (particle diameter less than about 45 μm) wollastonite particulate(a calcium silicate obtained from Peltz-Rowley Chemical Co.,Philadelphia, Pa.) was placed into the void space between theSiC/BN-coated fabric preform and the porous castable refractory boatuntil the level of the wollastonite was substantially flush with the topsurface of the preform. A thin layer of molten petroleum jelly

TABLE I Number of Average Max. Min. Sample Test Samples StrengthStrength Strength Condition Temp. Tested (MPa) (MPa) (MPa) Metallicconstituent Room temp. 8 461 ± 28 511 438 removed Metallic constituent1200° C. 10 488 ± 22 517 440 removed Metallic constituent 1300° C. 4 400± 12 412 386 removed Metallic constituent 1400° C. 4 340 ± 11 348 325removed Metallic constituent Room temp. 3 288 ± 21 302 264 removed andheat treated at 1200° C. in air for 24 h. Metallic constituent 1200° C.3 397 ± 9  404 387 removed and heat treated at 1200° C. in air for 24 h.Metallic constituent Room temp. 3 265 ± 12 275 253 removed and heattreated at 1200° C. in air for 100 h. Metallic constituent 1200° C. 3401 ± 28 433 319 removed and heat treated at 1200° C. in air for 100 h.

(VASELINE™, Cheesebrough-Ponds, Inc., Greenwich, Conn.) was firstapplied to the top surface of the SiC/BN-coated fabric preform and thencovered with nickel oxide (NiO) powder substantially in accordance ofthe methods of Example 2.

The porous castable refractory boat, having stilts at its corners, wasplaced into a resistance heated air atmosphere furnace and heated toabout 700° C. at a rate of about 400° C. per hour. A parent metal,comprising by weight about 7.5-9.5% Si, 3.0-4.0% Cu, ≦2.9% Zn, 0.2-0.3%Mg, ≦1.5% Fe, ≦0.5% Mn, ≦0.35% Sn, and the balance aluminum and weighingabout 420 grams, was also placed in a refractory container in theresistance heated air atmosphere furnace and heated to about 700° C.When parent metal was molten, the furnace door was opened and the parentmetal was poured into the heated porous castable refractory boat andonto the NiO powder coated preform, thereby covering the surface of theSiC/BN-coated fabric preform. Wollastonite powder was then placed ontothe surface of the molten parent metal within the porous boat tosubstantially completely cover the surface of the molten parent metal,thereby forming a lay-up. Then the furnace and its contents comprisingthe lay-up were heated to about 1000° C. in about an hour. After about60 hours at about 1000° C., the furnace and its contents were cooled toabout 700° C. in about 2 hours. At about 700° C., the lay-up was removedfrom the furnace and residual molten parent metal was decanted from theporous castable refractory boat. The refractory boat was rapidly brokenaway from the formed composite, and the formed composite was buried insilica sand to permit the composite to cool to about room temperature.At about room temperature, the composite was removed from the silicasand and it was observed that an oxidation reaction product comprisingalumina had grown into and substantially completely embedded theSiC/BN-coated fabric preform, thereby forming the ceramic matrixcomposite body having a plurality of fabric layers of 12 HSW ceramicgrade NICALON® fibers silicon carbide as a reinforcement. The ceramicmatrix also comprised some residual unreacted parent metal. The siliconcarbide fiber reinforced alumina composite body was then cut into barsmeasuring about 2.4 inches (60 mm) long by about 0.2 inch (6 mm) wide byabout 0.11 inch (3 mm) thick in preparation for the removal of at leasta portion of the metallic constituent of the formed fiber reinforcedceramic composite body.

Sample B

A graphite foil box having an inner cavity measuring about 4.0 inches(102 mm) long by about 4.0 inches (102 mm) wide by about 3.0 inches (96mm) deep was made from a piece of graphite foil (GRAFOIL™, UnionCarbide, Carbon Products Division, Cleveland, Ohio) measuring about 10.0inches (254 mm) long by about 10.0 inches (254 mm) wide by about 0.015inch (0.38 mm) thick. Four parallel cuts, 3.0 inches (76 mm) from theside and about 3.0 inches (76 mm) long were made into the graphite foil.The cut graphite foil was then folded and stapled to form the graphitefoil box.

A parent metal, comprising by weight about 3 percent strontium and thebalance aluminum and measuring about 4.0 inches (102 mm) long by about4.0 inches (102 mm) wide by about 1.0 inch (25 mm) thick was coated onone side thereof measuring about 4.0 inches (102 mm) long by about 4.0inches (102 mm) wide with a slurry comprising by weight about 90% −325mesh (particle size less than about 45 μm) aluminum alloy powder and thebalance ethanol. The −325 mesh aluminum alloy powder was nominallycomprised by weight of about 7.5-9.5% Si, 3.0-4.0% Cu, ≦2.9% Zn,0.2-0.3% Mg, ≦1.5% Fe, ≦0.5% Mn, ≦0.35% Sn, and the balance aluminum.The aluminum alloy powder-coated parent metal was then placed into thegraphite foil box such that the uncoated surfaces of the parent metalcontacted the inner surfaces of the graphite foil box.

A fabric preform measuring about 4.0 inches (102 mm) long by about 4.0inches (102 mm) wide by about 0.06 inch (1.6 mm) thick was made withinthe graphite foil box and on the aluminum alloy powder coated surface ofthe parent metal by stacking four layers of 12 harness satin weave (HSW)silicon carbide fabric (ceramic grade NICALON® silicon carbide fibrousmaterial obtained from Dow Corning Corporation, Midland, Mich.) onto theparent metal. About 0.5 inch (13 mm) of a 500 grit (average particlediameter of about 17 μm) alumina powder (El ALUNDUM™, Norton Company,Worcester, Mass.) was poured over the 12 HSW fabric preform and leveled.The sides of the graphite foil box that extended beyond the level of thealumina powder covering the 12 HSW fabrics were folded over onto thealumina powder to form a lid for the graphite foil box.

A lay-up was formed in a graphite refractory container by placing andleveling about 0.5 inch (13 mm) of a 500 grit (average particle diameterof about 17 μm) alumina powder into the bottom of the graphiterefractory container. The graphite foil box and its contents comprisingthe aluminum alloy powder-coated parent metal and the 12 HSW siliconcarbide fabric preform were placed into the graphite refractorycontainer and onto a 500 grit (average particle diameter of about 17 μm)alumina. Additional 500 grit alumina was placed into the graphiterefractory container into the void defined by the inner surface of thegraphite refractory container and the outer surface of the graphite foilbox. The 500 grit (average particle diameter of about 17 μm) aluminapowder also covered the top lid of the graphite foil box and itscontents.

The lay-up comprising the graphite refractory container and its contentswas placed into a retort lined resistance heat furnace and the retortdoor was closed. The furnace and its contents were heated to about 100°C. at a rate of about 300° C. per hour. At about 100° C., the retort wasevacuated to about 30.0 inches (762 mm) mercury (Hg) vacuum andmaintained at about 30.0 inches (762 mm) Hg vacuum to about 150° C. Atabout 150° C., nitrogen was introduced into the retort at a flow rate ofabout 4 liters per minute. The furnace and its contents were then heatedto about 900° C. at about 300° C. per hour. After about 200 hours atabout 900° C., the furnace and its contents were cooled to about roomtemperature at a rate of about 300° C. per hour. At about roomtemperature, the retort door was opened and the lay-up was removed. Thelay-up was disassembled, the preform was removed from within thegraphite foil box, and it was observed that an oxidation reactionproduct comprising aluminum nitride had grown into and substantiallycompletely embedded the silicon carbide fabric preform thereby forming aceramic matrix composite body reinforced with a plurality of fabriclayers of 12 HSW ceramic grade NICALON® silicon carbide asreinforcement. The ceramic matrix also comprised a metallic constituentcomprising residual unreacted parent metal.

Table II contains a summary of the parameters used to practice the metalremoval step of the instant invention on Samples A and B. Specifically,Table II contains the dimensions of the sample, the filler material usedfor metal removal, the infiltration enhancer precursor, the processingtemperature, the processing time at the processing temperature, and theprocessing atmosphere.

FIG. 6 shows a cross-sectional schematic of the setup used in thisseries of tests to remove the metallic constituent from Samples A and B.

After the formation of the silicon carbide fiber reinforced aluminacomposite body of Sample A had been achieved, the metal removal processwas effected. Specifically, a filler material mixture was formed,comprising by weight about 90 percent filler, which comprised 1000 grit(average particle diameter of about 5 μm) Al₂O₃ (E67 tabular alumina,Norton Co., Worcester, Mass.) and about 10 percent by weight −325 mesh(particle diameter less than about 45 μm) magnesium powder (AESAR™,Johnson Matthey, Seabrook, N.H.). The filler material mixture was mixedin a plastic jar on a rotating jar mill for about an hour.

A graphite foil box having an inner cavity measuring about 3 inches (76mm) long by about 3 inches (76 mm) wide and about 2.5 inches (64 mm)deep was made from graphite foil (PERMA FOIL, TT America, Portland,Oreg.). The graphite foil box was made from a piece of graphite foil,measuring about 8 inches (203 mm) long by about 8 inches (203 mm) wideby about 0.15 inches (4 mm) thick. Four parallel cuts about 2.5 inches(64 mm) from the side and about 2.5 inches (64 mm) long, were made intothe graphite foil. The graphite foil was then folded into a graphitefoil box

TABLE II Processing Infiltration Time At Filler Material EnhancerProcessing Processing Sample ID Composite¹ Geometry For Metal RemovalPrecursor Temperature Temperature Atm. A SiC_(f)/Al₂O₃ bar 1000 gritAl₂O₃ ² 10% −325 mesh Mg³ 850° C. 10 h N₂ B SiC_(f)/AlN irregular 1000grit Al₂O₃ 10% −325 mesh Mg³ 750° C. 10 h N₂ ¹SiC fiber reinforcedcomposite ²E-67 alumina, Norton Co., Worcester, NA. ³AESAR ® , JohnsonMatthey Corporation, Seabrook, New Hampshire

and stapled together. Metal was removed from Sample A by first pouringabout 0.5 inch (13 mm) of the mixture of filler material and magnesiumpowder into one of the graphite foil boxes. The filler material mixturewas levelled and hand tapped until smooth. A bar of the silicon carbidefiber reinforced alumina composite of Sample A, and measuring about 1.7inches (43.8 mm) long by about 0.25 inch (6.3 mm) wide by about 0.2 inch(4.5 mm) thick was placed onto the filler material mixture within thegraphite foil box and covered with another about 0.5 inch (13 mm) of thefiller material mixture which was again levelled and hand tapped untilsmooth.

The graphite foil box containing Sample A was then placed into agraphite refractory container having inner dimensions of about 9 inches(229 mm) long by about 9 inches (229 mm) wide by about 5 inches (127 mm)deep and having a wall thickness of about 0.5 inch (13 mm). The graphiterefractory container and its contents were then placed into a controlledatmosphere resistance heated furnace, the furnace door was closed andthe furnace was evacuated to about 30 inches (762 mm) Hg. After about 15hours at about 30 inches (762 mm) of mercury vacuum, the vacuum was shutoff and nitrogen gas was introduced into the furnace chamber at a flowrate of about 1 liter/minute. The operating pressure of the chamber wasabout 16.7 pounds per square inch (1.2 kg/cm ) with a nitrogen flow rateof about 1 liter/minute. The furnace was heated to about 850° C. atabout 200° C. per hour. After about 10 hours at about 850° C., the powerto the furnace was interrupted and the graphite refractory container andits contents were allowed to cool within the furnace to about roomtemperature. Once at room temperature, the graphite refractory containerand its contents were removed and the lay-up for Sample A wasdisassembled to reveal that the metallic constituent comprising analuminum alloy in the silicon carbide fiber reinforced alumina compositehad been drawn out from the composite body during the metal removalprocess.

The setup for the removal of the metallic constituent from Sample B wassubstantially the same as that described for Sample A of this Exampleand is schematically illustrated in FIG. 6. The nitrogen flow rate toeffect removal of the metallic constituent from Sample B was about twoliters per minute. The controlled atmosphere furnace was heated to aboutthe processing temperature of about 750° C. at a rate of about 200° C.per hour, held at about the processing temperature for about 10 hours.After about 10 hours at the processing temperature, at least a portionof the metallic constituent was removed from within the ceramic matrixcomposite body. Specifically, the metallic constituent spontaneouslyinfiltrated the filler material mixture comprising substantially a 1000grit (average particle diameter of about 5 μm) alumina and a −325 meshmagnesium infiltration enhancer precursor. The furnace and its contentswere cooled to about room temperature. At about room temperature, thesetup was removed from the furnace, disassembled, and weight loss due tothe removal of the metallic constituent from Sample B was noted.

EXAMPLE 4

The following Example demonstrates that fiber reinforced ceramiccomposite bodies formed by the method of the present invention maintainsubstantially their room temperature fracture toughness at elevatedtemperatures. A series of fiber preforms were made substantially inaccordance with the methods described in Example 2, except that thefirst layer and eighth layer of the fabric preform comprised 12 harnesssatin weave (12 HSW) fabric instead of 8 harness satin weave (8 HSW)fabric and the temperature of the methyltrichlorosilane (MTS) bath usedduring the formation of silicon carbide coatings was maintained at about18° C. instead of about 25° C. The lay-up for the growth of the fiberreinforced ceramic composite body included an alumina-bondedwollastonite shell fabricated substantially in accordance with themethods described in Example 2, and the composite growth process wassubstantially the same as that described in Example 2. The resultantceramic matrix composite bodies were subjected to a metal removaltreatment substantially the same as that described in Example 2. Thesamples were subsequently machined to form mechanical test samples whichwere used to determine both the flexural strength and the fracturetoughness of the fiber reinforced ceramic composite bodies both as afunction of test temperature.

Table III summarizes the results of these tests. The methods formeasurement of the flexural strength was substantially in accordancewith the methods described in Example 2. The method of Munz, Shannon andBubsey (International Journal of Fracture, Vol. 16 (1980) R137-R141) wasused to determine the fracture toughness of the silicon carbide fiberreinforced ceramic composite bodies. The fracture toughness wascalculated from the maximum load of Chevron notch specimens in fourpoint loading. Specifically, the geometry of each Chevron notch specimenwas about 1.8 to 2.2 inches (45 to 55 mm) long, about 0.12 inch (3 mm)wide and about 0.15 inch (3.75 mm) high. A Chevron notch was cut in eachspecimen with a diamond saw to permit the propagation of a crackstarting at the notch and traveling through the sample. The Chevronnotched specimens, having the apex of the Chevron notch pointingdownward, were placed into a fixture within a Universal test machine.The notch of the Chevron notch specimen, was placed between two pinsabout 1.6 inches (40 mm) apart and about 0.79 inch (20 mm) from eachpin. The top side of the Chevron notch specimen was contacted by twopins about 0.79 inch (20 mm) apart and about 0.39 inch (10 mm) from thenotch. The maximum load measurements were made with a Syntec ModelCITS-2000/6 universal testing machine (System Integration TechnologyIncorporated, Stoughton, Mass.). A crosshead speed of 0.02 inches/minute(0.58 millimeters/minute) was used. The load cell of the universaltesting machine was interfaced to a computer data acquisition system.The Chevron notch sample geometry and maximum load were used tocalculate the fracture toughness of the material. Several samples wereused to determine an average fracture toughness for a given group ofparameters (e.g., temperature, fiber reinforced ceramic composite body,etc.)

Table III summarizes the results of the measurements of the averageflexural strength, the maximum flexural strength and the averagefracture toughness all as a function of temperature, for Samples D, Eand F, which were subjected to the metal removal process. Moreover, thefracture toughness of an “as-grown” Sample C (e.g., without any residualmetallic constituent removed) is compared to a treated Sample D (i.e.,metallic constituent removed). The data in Table III shows that thefracture toughness of a fiber reinforced ceramic composite body with itsmetallic constituent substantially completely removed is notsignificantly diminished at elevated temperatures. In addition, thefracture toughness of a sample which is subjected to the metal removalprocess does not appear to vary significantly from the fracturetoughness of an untreated composite body.

EXAMPLE 5

The following Example demonstrates that fiber reinforced ceramiccomposite bodies exhibiting excellent fracture toughness can be producedby (1) coating a fabric preform with coatings comprising silicon carbide(SiC)/boron nitride (BN); (2) growing an oxidation reaction product by areaction of a parent metal with an oxidant which embeds theSiC/BN-coated fabric preform and (3) removing at least some of themetallic constituent from the grown fiber reinforced ceramic compositebody.

A ceramic grade NICALON® silicon carbide fiber reinforced aluminacomposite body plate measuring substantially the same as that in Example2 was formed substantially in accordance with the method of Example 2.Specifically, the fabric preform lay-up, the formation of both the boronnitride and silicon carbide coatings, the growth of the aluminaoxidation reaction product embedding the SiC/BN-coated fabric preformand the removal of the metallic constituent from the fiber reinforcedceramic body were performed substantially in accordance with the methodof Example 2.

TABLE III Average Maximum Average Flexural Flexural Fracture SampleSample Test Strength Strength Toughness ID Condition Temp. (MPa) (MPa)(MPa-m^(1/2) ) C As Grown RT — — 19 ± 1 D Metallic RT 450 (31)* 563 21 ±1 constituent removed E Metallic 1000° C. 400 (7)*  432 23 ± 1constituent removed F Metallic 1200° C. 350 (14)* 406 18 ± 1 constituentremoved *The number in parentheses indicates the number of sample test.

The fracture toughness of the fiber reinforced ceramic composite bodywas measured substantially in accordance with the method of Example 4,except that specimen size used to determine the toughness measured fromabout 1.0 to about 1.2 inches (25 to 30 mm) long, about 0.15 inch (3.75mm) high and about 0.12 inch (3 mm) wide. The apex of the Chevron notchpointed up within the universal test machine. The notch of the specimenwas placed between two pins about 0.39 inch (10 mm) apart and about 0.2inch (5 mm) from each pin. The top side of the specimen was contacted bytwo pins about 0.79 inch (20 mm) apart and about 0.39 inch (10 mm) fromthe notch. Three specimens were tested to determine an average fracturetoughness for a specific test temperature.

The fracture toughness of the fiber reinforced ceramic composite body ofthis Example was measured at about room temperature, at about 1200° C.and at about 1300° C. These values were about 35.3±1 MPa-m^(½), 19.6±1MPa-m^(½) and 18.7±1 MPa-m^(½), respectively.

EXAMPLE 6

The following Example demonstrates the intrinsic strength of the ceramicmatrix of a fiber reinforced ceramic composite body.

A ceramic grade NICALON® silicon carbide fiber reinforced aluminacomposite was formed substantially in accordance with the methods ofExample 2. Specifically, the fabric preform lay-up, the formation ofboth the boron nitride and silicon carbide coatings, the growth of thealumina oxidation reaction product embedding the SiC/BN-coated fiber andthe removal of the metallic constituent from the fiber reinforcedceramic body were performed substantially in accordance with the methodof Example 2.

The intrinsic strength of the matrix was measured at about roomtemperature with the short beam method according to ASTM method D2344-84 entitled “Standard Test Method for Apparent Interlaminar ShearStrength of Parallel Fiber Composite By Short-Beam Method.”

The mechanical test specimens were machined to overall dimensions ofabout 1 inch (25 mm) in length by about 0.16 inch (4 mm) in width byabout 0.16 inch (4 mm) in thickness. Furthermore, the orientation of themechanical test specimens were such that all the fibers wereperpendicular to the thickness dimension, i.e., none of the fiberstraversed the thickness dimension.

This test was specifically designed to measure the strength, and inparticular, the shear strength, of the matrix material between twoadjacent layers of the eight total layers of HSW fabric.

A three-point flexural test was used. The thickness and width of thetest bars was measured with a precision of about 390 microinch (0.01mm). The test bars were subjected to a stress which was applied at threepoints by two lower span bearing points and one upper span bearingpoint. The lower span bearing points were about 0.67 inch (17 mm) apartand the upper load point was centered over the lower span so that theload was applied substantially symmetrically on the test bar. Theflexural strength measurements were made with a Syntec Model No.CITS-2000/6 universal testing machine (System Integration Technology,Inc., Stoughton, Mass.) having a 500 pound (2225 N) full-scaledeflection load cell. A computer data acquisition system was connectedto the measuring unit and strain gauges in the load cell recorded thetest responses. The cross-head speed during testing was about 0.05 inchper minute (1.3 mm per minute).

The interlaminar shear strength was found to be about 62 MPa.

EXAMPLE 7

This Example characterizes the tensile strength of a fiber reinforcedceramic composite body and shows the gradual and progressive failure ofsuch a body as opposed to the sudden and catastrophic failure typical ofmost ceramic or ceramic composite bodies.

A ceramic grade NICALON® silicon carbide fiber reinforced alumina matrixcomposite was formed substantially in accordance with the method ofExample 2. Specifically, the fabric preform lay-up, the formation ofboth the boron nitride and silicon carbide coatings, the growth of thealumina oxidation reaction product embedding the SiC/BN-coated fiber andthe removal of the metallic constituent from the fiber reinforcedceramic body were performed substantially in accordance with the methodof Example 2.

The tensile strength of the fiber reinforced ceramic composite body wasmeasured using the procedures described in ASTM designations A 370 and E8M-88.

FIG. 7 shows the approximate shape of the test specimen which wasmachined using diamond grinding with the longitudinal axis of the testspecimen parallel to either the length or width dimension of the fiberpreform. The tensile test specimen measured overall about 6 inches (152mm) long by about 0.5 inch (13 mm) wide by about 0.12 inch (3 mm) thick.The gage section measured about 0.75 inch (19 mm) long by about 0.35inch (9 mm) wide. The test was performed using an MTS Model 810universal testing machine (MTS Systems Corp., Eden Prarie, Minn.)operated at a crosshead speed of about 0.25 mm per minute. The samplestrain was monitored with an MTS Model 632-11B-20 clip-on extensometer(MTS Systems Corp.).

At room temperature, the average tensile strength for 14 samples wasabout 331 MPa with a standard deviation of about 22 MPa. The Young'sModulus, as measured by the ratio of stress to strain in the linearportion of the stress-strain curve, averaged about 162 GPa and theaverage strain-to-failure was about 0.645 percent.

FIG. 8 shows a typical stress-strain curve for a fiber reinforcedceramic composite body made substantially by the method of Example 2.The stress-strain curve begins to deviate from linearity at a stress ofabout 50-60 MPa, which deviation indicates the onset of matrixmicrocracking and pull-out of the reinforcing fibers from thesurrounding matrix material.

FIG. 9 is a scanning electron micrograph taken at about 50×magnification of a fracture surface which has been exposed as a resultof a room temperature tensile test. Segments of the reinforcing fiberswhich have been partially pulled out of the surrounding matrix materialare clearly visible.

EXAMPLE 8

This Example demonstrates that fiber reinforced ceramic matrixcomposites produced according to the method of the present inventionretain almost all of their ambient temperature strength at elevatedtemperatures, even after repeated thermal cycling.

A fabric preform 103 was made by stacking a plurality of layers of 8harness satin weave (8 HSW) fabric and 12 harness satin weave (12 HSW)fabric made from NICALON® silicon carbide fiber (ceramic grade, obtainedfrom Dow Corning Corp., Midland, Mich.) on top of each other. Thenomenclature describing the orientations of the fabrics is substantiallythe same as that used in Example 2 and depicted in FIGS. 5a, 5 b and 5c.

The fabric preform of the present Example was made by stacking thelayers of HSW fabric in the following sequence:

A first fabric layer comprising an 8 HSW fabric was rotated about 90° inthe counterclockwise direction from the as-is position about an axis 93perpendicular to the plane of the fabric and was placed on a supportingsurface to start the fabric preform;

A second fabric layer comprising an 8 HSW fabric was placed on the firstfabric layer in the as-is position so that the edges of the secondfabric layer were substantially aligned with the edges of the firstfabric layer;

A third fabric layer comprising a 12 HSW fabric was rotated about 90° inthe counterclockwise direction from the as-is position about an axis 93perpendicular to the plane of the fabric and was placed on the secondfabric layer so that the edges of the third fabric layer weresubstantially aligned with the edges of the second fabric layer;

A fourth fabric layer comprising a 12 HSW fabric was placed on the thirdfabric layer in the as-is position so that the edges of the fourthfabric layer were substantially aligned with the edges of the thirdfabric layer;

A fifth fabric layer comprising a 12 HSW fabric was rotated about 90° inthe counterclockwise direction from the as-is position about an axis 93perpendicular to the plane of the fabric and was placed on the fourthfabric layer so that the edges of the fifth fabric layer weresubstantially aligned with the edges of the fourth fabric layer;

A sixth fabric layer comprising an 8 HSW fabric was placed on the fifthfabric layer in the as-is position so that the edges of the sixth fabriclayer were substantially aligned with the edges of the fifth fabriclayer;

A seventh fabric layer comprising an 8 HSW fabric was rotated about 90°in the counterclockwise direction from the as-is position about an axis93 perpendicular to the plane of the fabric and was placed on the sixthfabric layer so that the edges of the seventh fabric layer weresubstantially aligned were substantially aligned with the edges of thesixth fabric layer, thus completing the rectangular fabric preform whichmeasured about 7 inches (178 mm) in length by about 5 inches (127 mm) inwidth.

The fabric preform was clamped in substantially the same fixture as wasdescribed in Example 2 and depicted in FIG. 5e. The preform containmentfixture 108 containing the fabric preform was placed into a reactorchamber of a refractory alloy steel chemical vapor infiltrationapparatus having a graphite tube liner and having overall dimensions ofabout 8 feet (2.4 meters) in length by about 15.5 inches (394 mm) ininside diameter. The warp yarns of the first and seventh layers of thefabric preform were perpendicular to the gas flow direction within thechamber as well as being perpendicular to the longitudinal axis of thereactor chamber. The reactor chamber was closed and evacuated to lessthan about 0.04 inch (1 mm) of mercury (Hg). The reactor chamber wasthen heated to a temperature of about 820° C. Argon gas was flowed intothe annulus region between the graphite liner and the steel reactor wallat a rate of about 850 standard cubic centimeters per minute (sccm).When the temperature within the reactor chamber reached about 820° C., agas mixture comprising borontrichloride (BCl₃) flowing at about 700 sccmat a temperature of about 60° C. and ammonia (NH₃) flowing at about 180sccm was introduced into the reactor chamber while maintaining a totaloperating pressure of about 0.5 torr. After about 7 hours at atemperature of about 820° C., the gas mixture flowing into the reactorchamber was interrupted, the power to the furnace heating the reactorchamber was interrupted and the furnace and its contents were naturallycooled. At a temperature below about 200° C., the reactor chamber doorwas opened and the graphite containment fixture was removed, cooled anddisassembled to reveal that the fibers of the fabric layers of thefabric preform were coated and that the fabric layers comprising thefabric preform were bonded together by a boron nitride coating. Theboron nitride coating had a thickness of about 0.48 micron.

The boron nitride coated fabric preform was then stored in a vacuumdesiccator until it was ready to be put back into the chemical vaporinfiltration apparatus for additional coating.

For the application of this subsequent coating, the boron nitride coatedand bonded fabric preform was placed back into the reactor chamber ofthe chemical vapor infiltration apparatus. In this instance, however,the warp yarns of the first and seventh layers of the fabric preformwere parallel to the gas flow direction within the chamber, as well asbeing parallel to the longitudinal axis of the reactor chamber. Thereactor chamber was closed and evacuated to about less than about 1torr. Hydrogen gas was introduced into the reactor chamber at a flowrate of about 5000 standard cubic centimeters per minute (sccm). Thereactor chamber was then heated to a temperature of about 935° C.Nitrogen gas was flowed through the annulus region at a rate of about850 sccm. Once the temperature of the contents of the reactor chamberhad substantially completely stabilized at about 935° C., about 1500sccm of hydrogen were diverted away from direct entry into the reactorchamber and were first bubbled through a bath of methyltrichlorosilane(MTS) maintained at a temperature of about 45° C. before entering thereactor chamber. After about 20 hours at a temperature of about 935° C.,the power to the furnace heating the reactor chamber was interrupted andthe about 1500 sccm of hydrogen that was being directed through the MTSbath was again permitted to flow directly into the reactor chamber tore-establish a direct hydrogen gas flow rate of about 5000 sccm into thereactor chamber. After the reactor chamber had cooled substantially, thehydrogen flow rate was interrupted and the furnace and its contents wereevacuated to less than 1 torr. The pressure within the reactor chamberwas then brought back up to about atmospheric pressure with argon gas.After the reactor chamber had cooled to a temperature below about 200°C., the argon gas flow rate was interrupted and the reactor chamber doorwas opened. The graphite containment fixture was removed, cooled anddisassembled to reveal that the boron nitride bonded fabric preform hadbeen coated with a second layer of silicon carbide thereby forming asilicon carbide (SiC)/boron nitride (BN)-coated fabric preform. Thesilicon carbide had a thickness of about 1.9 microns.

Growth of an alumina oxidation reaction product through the siliconcarbide/boron nitride-coated fabric preform was then carried out insubstantially the same manner as was described in Example 2 to form afiber reinforced ceramic composite body comprising a ceramic matrixcomprising an aluminum oxide oxidation reaction product and a metalliccomponent in comprising some residual unreacted parent metal, with saidceramic matrix reinforced by the silicon carbide/boron nitride coatedNICALON® silicon carbide fibers (ceramic grade). Substantially completegrowth of the ceramic matrix only required about 72 hours, however.

Once the ceramic composite body had been manufactured, at least aportion of the metallic constituent comprising the ceramic matrix wasremoved. This metal removal process was performed in substantially thesame manner as was described in Example 2.

Tensile test specimens were machined from the fiber reinforced ceramiccomposite body and tested in substantially the same manner as describedin Example 7. Heating was provided by positioning a resistance heatedair atmosphere furnace in the testing zone of the test machine. Thesamples were tested in air at ambient as well as at elevatedtemperatures of about 1100° C., 1200° C., and about 1370° C. As shown inFIG. 10, the tensile strength at these temperatures was about 260, 250,260, and about 230 MPa, respectively. Thus, these data show that thefiber reinforced ceramic composite material retains substantially all ofits ambient temperature strength up to a temperature of about 1200° C.,and almost all of its ambient temperature strength at a temperature ofabout 1370° C.

Next, the effect of repeated thermal cycling on the material's tensilestrength was assessed.

First, a ceramic grade NICALON® silicon carbide reinforced aluminamatrix composite was produced substantially in accordance with theprocedure described in Example 2 and likewise subjected to the metalremoval process of Example 2. Unlike the procedure of Example 2,however, during the growth of oxidation reaction product into thepreform, the approximately 950° C. process temperature was maintainedfor about 100 hours instead of about 90 hours. Moreover, the thicknessof the silicon carbide coating deposited onto the boron nitride coatedNICALON® silicon carbide fibers during chemical vapor infiltration wasabout 2.0 microns.

Substantially rectangular tensile test specimens were diamond machinedfrom the composite tile such that the length dimension of the testspecimen was oriented parallel to the length or width dimension of thecomposite tile.

About half of the specimens were given a rapid thermal cycling treatmentbefore tensile testing; the others were tested “as is”. Specifically,the thermal cycling comprised subjecting each composite test specimen toabout 150 thermal cycles, each thermal cycle comprising heating a testspecimen from a starting temperature to a temperature of about 1200° C.in an argon atmosphere at a rate of about 40° C. per minute, holding ata temperature of about 1200° C. for about 2 minutes, and cooling back tothe starting temperature at a rate of about 10° C. per minute. Thestarting temperature corresponded to the final testing temperature. Thetwo sets of tensile test specimens were then tested in substantially thesame manner as was described in the preceding Example at about roomtemperature and temperatures of about 1000° F. (538°) about 1500° F.(816° C.) and at about 2000° F. (1093° C.).

FIG. 11 shows the tensile strength as a function of test temperature forthe two sets of composite test specimens. The data show that thethermally cycled composite test specimen experienced little loss intensile strength compared to their counterparts which were not thermallycycled. The significance of this result is that the thermal cyclingprovided an opportunity for chemical reaction between the fiber, thefiber coatings and the surrounding matrix constituents. The thermalcycling operation also provided an opportunity for cracking due tothermal expansion mismatch. The lack of significant strength reductionindicates that any microcracking induced by the thermal cycling wasconfined to the matrix material and, furthermore, that the ability ofthe fibers to pull out of the matrix under the applied tensile load wasnot substantially affected by the thermal cycling. The different tensilestrength levels observed in comparing the data of FIG. 10 to that ofFIG. 11 may be attributable to variations in preform fabrication,specifically, such as the differences in the number of each type of HSWfabric (e.g., 12 HSW vs 8 HSW).

EXAMPLE 9

This Example demonstrates the high temperature mechanical performance ofa fiber reinforced ceramic composite body under an applied load over aprolonged period of time in an oxidizing atmosphere.

The fiber reinforced ceramic composite body described herein wasfabricated substantially in accordance with the methods outlined inExample 2. Specifically, the fabric preform lay-up, the formation ofboth the boron nitride and silicon carbide coatings, the growth of thealumina oxidation reaction product embedding the SiC/BN-coated fiber andthe removal of the metallic constituent from the fiber reinforcedceramic body were performed substantially in accordance with the methodof Example 2.

In Example 7, it was demonstrated that at room temperature (e.g., about20° C.) in a pure tensile test, a fiber reinforced ceramic matrixcomposite sample begins to deviate from linear stress/strain behavior atan applied stress of about 50-60 MPa, indicating that the matrix beginsto microcrack at approximately this stress level. These microcracks mayallow for oxygen in the surrounding atmosphere to find a path to theunderlying ceramic grade NICALON® silicon carbide fiber and/or its SiCand BN coatings. Accordingly, stress rupture tests were conducted atvarious elevated temperatures in air at applied stresses above this50-60 MPa microcracking threshold in order to evaluate the impact ofmatrix microcracking and subsequent oxygen ingress on the performance ofthe fiber reinforced ceramic composite body.

The stress rupture test specimen had substantially the same shape asthat depicted in FIG. 7, with the exception that shoulders were machinedinto each end of the test specimen so that the sample could be grippedby a collar in the test fixture rather than clamped. Mica powder wasused in the collar to cushion the contact zone between the collar andthe shoulder portions of the stress rupture test specimen. The testspecimen measured about 5.5 inches (140 mm) long overall by about 0.5inch (13 mm) wide by about 0.12 inch (3 mm) thick. The gage portion ofthe test specimen measured about 2 inches (51 mm) in length by about 0.2inches (5 mm) wide.

The tests comprised heating the samples to the desired test temperatureand loading each specimen in tension to a desired stress and maintainingsaid stress at said temperature. The applied stress was increased in astep-wise manner. The unit length change of the specimen within the gageportion of the overall test specimen was monitored with a Model 1102ZYGO™ helium-neon laser extensometer (Zygo Corp., Middlefield, Conn.).

The results of the stress rupture testing are presented for FIG. 12.

The particulars of the applied stress and the exposure times arepresented below.

Sample G

The test fixture, comprising the Sample G test specimen with collarsattached to each end, was loaded into a Model P-5 creep testing machine(SATEC Inc., Grove City, Pa.). A tensile stress of about 12.5megapascals was applied to the test specimen using dead loading. Aresistance heated air atmosphere furnace was positioned completelyaround the stress rupture test specimen and the furnace and the stressrupture sample contained within were heated from about room temperatureto a temperature of about 1000° C. over a period of about 2 hours.

After the furnace chamber and its contents had reached a temperature ofabout 1000° C., the stress applied to the sample was increased to about75 MPa. After maintaining an applied stress of about 75 MPa for about 70hours, the applied stress to the sample was increased to about 100 MPa.After about 15 hours at a stress of about 100 MPa, the sample broke. Thefurnace chamber and its contents were allowed to cool naturally backdown to about room temperature.

Sample H

The Sample H test fixture was placed into the creep testing machine atabout room temperature and the Sample H stress rupture test specimen washeated in the surrounding resistance heated air atmosphere furnace to atemperature of about 1000° C. over a period of about 3 hours under anapplied stress of about 5 MPa. At a temperature of about 1000° C., theapplied tensile stress on the sample was increased to about 70 MPa andthe temperature inside the furnace chamber was increased to about 1100°C. over a period of about 1 hour. After maintaining the sample intension at a stress of about 70 MPa at a temperature of about 1100° C.for about 210 hours, the applied stress was increased to about 83 MPa.After about an additional 6 hours, the stress was increased to about 85MPa. After maintaining an applied stress of about 85 MPa on the samplefor about 115 hours, the applied stress was increased to about 88 MPa.After maintaining an applied stress of 88 MPa for about 1.5 hours, thestress applied was increased to about 90 MPa. After maintaining anapplied tensile stress of about 90 MPa for about 3 hours, the appliedstress was increased to about 91 MPa. After maintaining an appliedstress of about 91 MPa for about 1.5 hours, the stress was furtherincreased to about 92 MPa. After maintaining an applied stress of about92 MPa for about 1.3 hours, the applied stress was increased to about 95MPa. After maintaining an applied stress of about 95 MPa on the samplefor about 115 hours, the applied stress was increased to about 96 MPa.After maintaining an applied stress of about 96 MPa for about 3 hours,the applied stress was increased to about 97 MPa. After maintaining anapplied stress of about 97 MPa for about 2 hours, the applied stress wasincreased to about 99 MPa. After maintaining an applied stress of about99 MPa for about 1.5 hours, the applied stress was increased to about100 MPa. After maintaining an applied stress of about 100 MPa for about60 hours, the sample broke. The furnace chamber and its contents werethereafter furnace cooled from a temperature of about 1100° C. down toabout room temperature.

The fractured sample was recovered from the test chamber and thefracture surface was examined in the scanning electron microscope. FIG.13 is an approximately 50× magnification scanning electron micrograph ofa portion of the fracture surface. Direct comparison of FIG. 13 with theprevious scanning electron micrograph of FIG. 9 shows much less fiberpull-out associated with this Sample H specimen than with the fracturesurface of the Example 7 tensile test specimen. This decrease in thedegree of fiber pull-out of the present stress rupture may suggestdegradation of the fiber and/or one or more of its coatings over the500+ hour duration of the stress rupture test. Conversely, the abilityof this fiber reinforced ceramic matrix composite body to survivesustained exposure of this duration at a temperature of about 1100° C.at a stress level sufficient to expose the reinforcing fibers and/ortheir coatings to atmospheric oxygen may suggest the operation of amechanism working to protect the NICALON® fibers from chemical reactionssuch as atmospheric oxidation.

FIGS. 14 and 15 are scanning electron micrographs taken at about 2500×,and 10,000× magnification of a diamond polished cross-section of theSample H stress rupture test specimen at a region very close to thefracture surface. Specifically, FIG. 14 shows a crack breaching at leastthe SiC coating, thus potentially exposing the NICALON® fiber and/or theBN debond coating to chemical reaction with reactant supplied fromoutside the fiber and its coatings. The higher magnification of thecrack region shown in FIG. 15 reveals the presence of a substance atleast partially filling the crack. Such a substance may comprise areaction product of one or both of the SiC and BN coatings and/or theNICALON® fiber itself. The presence of such a reaction product mayexplain the apparent degradation of the fiber pull-out mechanism as wellas the relative longevity of the material while under load at elevatedtemperature. Specifically, the at least partial re-filling of a matrixmicrocrack after such a crack forms may serve to reduce the access of,for example, atmospheric oxygen to the reinforcing fibers and theircoatings.

The matrix microcrack shown breaching an SiC coating in FIG. 15 isdifferent from the matrix microcrack shown in FIG. 14.

This particular micrograph appears to show that the substancesubstantially filling the crack in the SiC coating also substantiallycomprises the space between the SiC coating and the NICALON® fiber andthe space between the SiC coating and the alumina oxidation reactionproduct.

Sample I

Sample E was stress rupture tested at a temperature of about 1200° C.The sample was loaded into the test rig in substantially the same manneras was described for Sample G. A tensile stress of about 12.5 MPa wasapplied to the test specimen at about room temperature. The furnacechamber and its contents were then heated from about room temperature toa temperature of about 1200° C. over a period of about 3 hours. At atemperature of about 1200° C., the applied stress was increased to about66 MPa. After maintaining a temperature of about 1200° C. at an appliedstress of about 66 MPa for about 256 hours, the applied stress wasincreased to about 70 MPa. After maintaining an applied stress of about70 MPa at a temperature of about 1200° C. for about 216 hours, theapplied stress was increased to about 75 MPa. After maintaining anapplied stress of about 75 megapascals at a temperature of about 1200°C. for about 288 hours, the applied stress was increased to about 80MPa. After maintaining an applied stress of about 80 MPa at atemperature of about 1200° C. for about 242 hours, the applied stresswas increased to about 87 MPa. After about 1 hour at an applied stressof about 87 MPa at a temperature of about 1200° C., the sample broke.

Concurrent with the stress rupture test, the strain of the stressrupture test specimens was monitored in the gage portion of the testspecimen using the previously identified laser extensometer to helpassess the creep behavior of the fiber reinforced ceramic matrixcomposite test specimen. Specifically, the first portion of the stressrupture for Sample H was repeated. Instead of testing the sample tofailure, however, the temperature was decreased from about 1100° C. backdown to about room temperature after about 210 hours at about 1100° C.under the approximately 70 MPa applied tensile stress. FIG. 15 shows thecumulative percent strain in the gage portion of the Sample H testspecimen resulting from this creep test. The significance of FIG. 15 isthat during the course of this approximately 210 hour creep test, SampleH shows essentially no change in elongation, indicating substantially noplastic deformation of the sample. Accordingly, no creep deformation ofSample H occurred under the described test conditions.

Similarly, no creep deformation was observed in the Sample I materialwhich was stress rupture tested at a temperature of about 1200° C. underan applied load of about 70 MPa for about 216 hours. In contrast, it hasbeen demonstrated in the art that creep deformation occurs in ceramicgrade NICALON® silicon carbide fibers at about 1200° C. Accordingly, thepresent results suggest that the present particular disposition of thereinforcing fibers in the applied coatings and the surrounding matrixmaterial may provide enhanced creep resistance to the present fiberreinforced ceramic matrix composite system.

Furthermore, the present results may suggest that the particulardisposition of the reinforcing fibers in the present composite bodyprovides protection to said fibers from degradation (e.g., chemicalattack) such as from atmospheric gases (e.g., oxygen and nitrogen) atelevated temperatures. Specifically, one additional stress rupture testwas conducted on the fiber reinforced ceramic composite material of thepresent Example. The sample was tested in substantially the same manneras was Sample I except that after heating to a temperature of about1200° C., the applied tensile stress was increased from about 12.5 MPato about 80 MPa. After about 1000 hours in air at a temperature of about1200° C. and at a stress of about 80 MPa (as shown in FIG. 12), thesample broke.

EXAMPLE 10

This Example demonstrates the fabrication of a ceramic grade NICALON®silicon carbide fiber reinforced alumina matrix composite, wherein theNICALON® fibers are first CVD coated with dual boron nitride/siliconcarbide coatings applied in alternating layers starting with boronnitride.

A fabric preform was made by stacking 8 layers of 12 harness satin weave(12 HSW) fabric made from NICALON® silicon carbide fiber (ceramic grade,obtained from Dow Corning Corp., Midland, Mich.) on top of each othersubstantially in accordance with the procedure described for Sample A ofExample 3.

The fabric preform comprising the 8 layers of 12 HSW NICALON® siliconcarbide fabric were then placed into the graphite preform containmentfixture 108 described in Example 2 and depicted in FIG. 5e insubstantially the same manner as was described in Example 2. The preformcontainment fixture containing the fabric preform was then placed intothe reactor chamber of a chemical vapor infiltration apparatus having aninside diameter of about 4.5 inches (114 mm) and a length of about 18inches (457 mm). The warp yams of the eighth layer of the fabric preformwere parallel to the gas flow direction within the chamber as well asbeing parallel to the longitudinal axis of the reactor chamber. Thereactor chamber was closed and evacuated to less than about 0.6 torr.The reactor chamber was then heated to a temperature of about 800° C. bymeans of inductive heating. When the temperature within the reactorchamber reached about 800° C., as indicated by a thermocouple containedtherein, a gas mixture comprising ammonia (NH₃) flowing at about 400standard cubic centimeters per minute (sccm) and boron trichloride(BCl₃) flowing at about 200 sccm was introduced into the reactor chamberwhile maintaining a total operating pressure of about 0.6 torr. Afterabout 2 hours at a temperature of about 800° C., the gas mixture flowinginto the reactor chamber was interrupted, the power to the furnaceheating the reactor chamber was interrupted and the furnace and itscontents were naturally cooled. After sufficient cooling (e.g., lessthan about 200° C.), the reactor chamber door was opened and the preformcontainment fixture was removed, cooled and disassembled to reveal thatthe fibers of the fabric layers of the fabric preform were coated withboron nitride, and furthermore, that the fabric layers comprising thefabric preform were bonded together by the boron nitride coating. Theboron nitride coating thickness on the fibers was about 0.33 microns.

The boron nitride coated and bonded fabric preform was stored in avacuum desiccator pending subsequent coating.

Next, a silicon carbide coating was applied to the fibers of the fabricpreform.

The boron nitride coated and bonded fabric preform was placed back intothe reactor chamber of the above-described chemical vapor infiltrationapparatus. Because the fiber preform was self-bonding at this stage, thegraphite containment fixture was unnecessary. The orientation of thefabric preform, however, was substantially the same as that employed fordepositing the boron nitride coating onto the fibers in the previousdeposition reaction.

The reactor chamber door was closed and the reactor chamber and itscontents were evacuated to less than about 0.3 torr. The reactor chamberand its contents were then heated from about room temperature to atemperature of about 925° C. at a rate of about 50° C. per minute.Hydrogen gas was then introduced into the reactor chamber at a flow rateof about 750 standard cubic centimeters per minute (sccm). When thereactor chamber and its contents had equilibrated at a temperature ofabout 925° C., as indicated by a thermocouple contained therein,additional hydrogen flowing at a rate of about 750 sccm was bubbledthrough a liquid bath of methyltrichlorosilane (MTS) maintained at atemperature of about 21 ° C., after which this gas was introduced intothe reactor chamber. The pressure in the reactor chamber was stabilizedat about 11 torr. After maintaining these conditions of temperature,pressure and gas flow rate for about 3 hours, power to the resistanceheated furnace which heated the reactor chamber was interrupted and theabout 750 sccm of hydrogen that was being directed through the liquidMTS bath was diverted around the MTS bath and permitted to flow directlyinto the reactor chamber, thus establishing a direct hydrogen gas flowrate of about 1500 sccm into the reactor chamber. After the temperatureof the reactor chamber and its contents had dropped to about 800° C.,the resistance heated furnace was re-energized and the temperature ofthe reactor chamber and its contents was stabilized at about 800° C.

Another boron nitride coating was then deposited on the coated fiber.Specifically, the flow of hydrogen gas into the reactor was interruptedand the reactor chamber and its contents were then evacuated to lessthan about 0.3 torr. Ammonia (NH₃) and borontrichloride (BCl₃) gaseswere then introduced into the reactor chamber in substantially the samemanner as was described previously at an operating pressure of about 0.6torr so as to deposit a coating of boron nitride onto the coated fiberscomprising the fabric preform. After depositing boron nitride for about1.5 hours at a temperature of about 800° C. and at a pressure of about0.6 torr, the gas mixture flowing into the reactor chamber wasinterrupted. The temperature of the reactor chamber and its contents wasraised from about 800° C. back up to about 925° C. Hydrogen gas was thenreintroduced into the furnace chamber at a flow rate of about 750 sccm.

When the temperature of the reactor chamber and its contents hadstabilized at about 925° C., a final coating of silicon carbide wasdeposited onto the coated NICALON® silicon carbide fibers comprising thefabric preform.

Specifically, substantially the same procedure was employed indepositing this second silicon carbide coating as was employed indepositing the first silicon carbide coating described earlier, with theexception that the reactor chamber and its contents were maintained at atemperature of about 925° C. at an operating pressure of about 11 torrfor about 20 hours.

After depositing this second silicon carbide coating for about 20 hours,the power to the furnace heating the reactor chamber was interrupted andthe about 750 sccm of hydrogen which was bubbled through the liquid MTSbath was instead sent directly into the reactor chamber without firstbeing routed through the MTS bath. After the furnace chamber and itscontents had cooled down to about less than about 200° C., the flow ofhydrogen gas into the reactor chamber was interrupted and the reactorchamber was evacuated to less than about 0.3 torr. The pressure in thefurnace chamber was then returned to atmospheric pressure using argongas. When the furnace chamber had reached substantially atmosphericpressure, the chamber was opened and the coated fabric preform wasremoved from the reactor chamber.

An alumina oxidation reaction product was grown into the coated fiberpreform in substantially the same manner as was described for Sample Aof Example 3 to form a ceramic composite body comprising NICALON®silicon carbide fibers coated with, in order from interior to exterior,about 0.2 micron boron nitride, about 1.83 microns silicon carbide,about 0.2 micron boron nitride and about 1.93 microns silicon carbide asmeasured along the radius of the fiber cross-section, said coatedNICALON® fibers reinforcing a ceramic matrix, said ceramic matrixcomprising an alumina oxidation reaction product and a metallicconstituent comprising some residual unreacted parent metal.

At least a portion of the metallic component of the formed compositebody was then removed in substantially the same manner as described inExample 2.

Flexural strength test specimens were machined and strength tested atabout room temperature in substantially the same manner as was describedin Example 2. FIGS. 16a and 16 b are a scanning electron micrographs atabout 3500× magnification of a polished cross-section of the fracturesurface of the fiber reinforced ceramic composite test specimen. Inparticular, FIG. 16a shows a crack entering the outer silicon carbidelayer and exiting without going through the inner silicon carbide layer.Not all of the cracks displayed this behavior, however, as evidenced byFIG. 16b which shows a crack entering through both outer and innersilicon carbide coating layers and subsequently exiting through bothsilicon carbide layers.

Demonstration that a NICALON® fiber reinforced ceramic composite whoseNICALON® fibers have coated thereon double layers of boron nitride andsilicon carbide can fracture or debond between the inner and outersilicon carbide layers may suggest that those fibers where this behavioroccurs will be more resistant to chemical degradation from externalreactants at elevated temperatures because such fibers are stillprotected by one group of boron nitride and silicon carbide coatings.

EXAMPLE 11

This Example demonstrates that a coating of boron nitride followed by acoating of silicon carbide on a NICALON® fiber provide some protectionfrom oxidation at elevated temperatures. This Example also shows thatthe application of an additional set of boron nitride and siliconcarbide coatings supplied over the first set provide significantlygreater oxidation protection.

Thermogravimetric analyses were performed on Samples J, K, L and Mdescribed below. Each test comprised placing a sample having a mass ofseveral tens to several hundreds of milligrams into an alumina cruciblewhich in turn was placed into the test chamber of a Model SPA 409Netzsch microbalance (Netzsch Inc., Exton, Pa.). The chamber was sealedand substantially pure oxygen gas was introduced into the test chamberat a flow rate of about 200 standard cubic centimeters per minute(sccm). The temperature of the sample was then increased fromsubstantially room temperature to a temperature of about 1200° C. at arate of about 200° C. per hour. After maintaining a temperature of about1200° C. for about 24 hours, the temperature was decreased to about roomtemperature at a rate of about 200° C. per hour. The flow of thesubstantially pure oxygen gas was interrupted. The microbalancecontimuously monitored and recorded the mass of the test samplethroughout the duration of the test.

Sample J

Sample J comprised ceramic grade NICALON® fibers in the “as-received”condition.

Sample K

Sample K comprised ceramic grade NICALON® fibers which were coated withboron nitride substantially in accordance with the method described inExample 10.

Sample L

Sample L comprised ceramic grade NICALON® fibers which were coated witha layer of boron nitride and a layer of silicon carbide substantially asdescribed in Example 10.

Sample M

Sample M comprised a sample of the laminate which was deposited on thereactor wall in accordance with the procedure of Example 10. Thelaminate material specifically comprised, in succession, layers of boronnitride, silicon carbide, additional boron nitride and additionalsilicon carbide substantially as described in Example 10.

Table IV shows the percentage weight gain for each of the four samplesas a function of the initial sample weight (e.g., fiber weight plus theweight of any coatings), the weight only of initial NICALON® fiber, theweight only of the boron nitride coating, and the weights of the siliconcarbide and boron nitride coatings. For Sample M only the percentageweight increase in terms of the initial sample weight was measured.

The data show that coating a NICALON® fiber with both boron nitride andsilicon carbide substantially reduces the elevated temperature oxidationof the fiber in oxygenated environments, as evidenced by the weightincreases of 0.47 and 0.65 percent, respectively, compared to the weightincrease of 1.4 percent for an uncoated NICALON® fiber. Moreover, theTable appears to indicate that the best oxidation resistance (e.g., theleast amount of weight increase) may occur when a dual duplex coating ofboron nitride and silicon carbide (e.g., four layers in all) is applied.This result may suggest that this dual duplex coating could not onlyprotect a NICALON® fiber but could also protect the underlying boronnitride/silicon carbide coatings and in particular, the inner boronnitride debond coating.

Although only a few exemplary embodiments of the invention have beendescribed in detail above, those skilled in the art would readilyappreciate that the present invention embraces many combinations andvariations other than those exemplified.

EXAMPLE 12

This Example demonstrates, among other things, that the addition ofparticulates of silicon carbide between the fibric plies of fiber in afiber reinforced ceramic composite body can greatly reduce the extent ofmicrocracking in the ceramic

TABLE IV Weight Gains of Coated and Uncoated NICALON ™ Fibers and CVISiC/BN Coatings on Exposure to Air at 1200° C. for 24 Hours % WeightGain Based on Based on Based on Based on Sample Initial Initial CVI BNSiC/BN ID Description Sample Wt. Fiber Wt. Coating Wt. Coating Wt. JNICALON ™ Fiber 1.4 1.4 — — K BN Coated 13.63 15.39 119.53 — NICALON ™Fiber L SiC/BN Coated 0.47 2.0  14.81 0.65 NICALON ™ Fiber MSiC/BN/SiC/BN 0.08 — — — Coated NICALON ™ Fiber

matrix material which may occur between the plies during compositefabrication.

A graphite containment fixture containing a fabric preform was assembledin substantially the same manner as was described in Example 2, with thefollowing notable exceptions. First, the NICALON® silicon carbide fabric(certamic grade obtained from Dow Corning Corp., Midland, Mich.)entirely comprised 8-harness satin weave (8 HSW). Moreover, about 5.75grams of 220 grit 39 CRYSTOLON\ dry silicon carbide particulate (averageparticle size of about 66 microns, Norton Company, Worcester, Mass.) wasevenly applied to the top 6.75 inch (171 mm) face of each of fabricplies 2-6. Still further, the dry silicon carbide particulate was workedat least part way into the tows of fiber making up each fiber ply with abrush.

The fabric preform comprising the eight stacked layers of 8-harnesssatin weave (8 HSW) fabric containing the five layers of silicon carbideparticulate were then placed into a chemical-vapor infiltration (CVI)reactor and the fibers were coated with a first layer of boron nitride(BN) followed by layer of silicon carbide (SiC) substantially inaccordance with the method described in Example 2. As a result ofchemical-vapor infiltration, about 0.51 micron of boron nitride andabout 1.94 microns of silicon carbide, as calculated based upon preformweight gain, were deposited onto the reinforcement fibers andparticulates in the preform.

A ceramic matrix comprising aluminum oxide and some aluminum alloy metalwas then grown into the coated preform by means of the directed metaloxidation process described in Example 2 to form a fiber reinforcedceramic matrix composite body.

The formed ceramic composite body was then subjected to substantiallythe same metal removal process as described in Example 2 to remove atleast some of the metallic component of the ceramic composite body.

The ceramic composite body was then sectioned using a diamond saw,mounted in thermoplastic resin and polished using progressively finergrades of diamond paste to produce a sufficiently smooth surface foroptical examination. FIG. 17b is an approximately 50× magnificationoptical photomicrograph of this polished cross-section of this fiberreinforced ceramic matrix composite. Specifically, this figure showsparticulates of silicon carbide 306 embedded within an alumina ceramicmatrix material 302 located between and embedding adjacent plies offabric comprising woven tows of the reinforcement NICALON® fibers 304.FIG. 17a is also an approximately 50× optical photomicrograph of a fiberreinforced ceramic matrix composite which was produced in substantiallythe same manner as the composite material of the present Example withthe exception that no silicon carbide particulates were placed betweenthe fabric plies of NICALON® fiber. The absence in FIG. 17b of thecracks 300 shown in FIG. 17a are particularly noticeable andsignificant.

Thus, the present Example, among other things, demonstrates thataddition of a particulate material such as silicon carbide between theplies of NICALON® silicon carbide fabric can substantially reduce, ifnot completely eliminate, the phenomenon of microcracking in the matrixmaterial occupying the space between the plies of NICALON® fabric.

EXAMPLE 13

This Example demonstrates, among other things, that there is a preferredthickness for each of the boron nitride and silicon carbide coatingswhich are applied to a preform comprising silicon carbide reinforcementfibers if the optimal flexural strength is to be achieved. Moreparticularly, this Example demonstrates that for ceramic compositebodies having about 35 to about 36 volume percent reinforcement fibersin a matrix comprising predominantly aluminum oxide, the optimumthickness of boron nitride is somewhere between 0.20 micron and 0.41micron and the optimum thickness of silicon carbide is somewhere aboveabout 1.9 microns.

Sample N

A fabric preform was made in substantially the same manner as wasdescribed for Sample A of Example 3, except that all eight layers offabric comprising ceramic grade NICALON® silicon carbide fiber (obtainedfrom Dow Corning Corp., Midland, Mich.) comprised 12-harness satin weave(12 HSW) fabric. The fabric preform was then placed into a graphitecontainment fixture whose shape (but not necessarily size) wassubstantially as shown in FIG. 5e.

The graphite containment fixture containing the fabric preform was thenplaced into a reactor chamber of a chemical-vapor infiltration (CVI)apparatus having a outer diameter of about 4.5 inches (110 mm) and alength of about 18 inches (441 mm). The reactor was inductively heatedand oriented vertically such that the reactive gases were introduced atthe top of the reactor and exhausted at the base of the reactor.KANTHAL\ iron-chromium-aluminum alloy wires were used to suspend thegraphite containment fixture about 11.5 inches (282 mm) from the top ofthe reactor. The warp yarns of the eighth layer of the fabric preformwere parallel to the gas flow direction within the chamber. Thesubsequent processing was then substantially the same as that utilizedto deposit the first boron nitride coating of Example 10 with theexception that the reactor was maintained at the coating temperature ofabout 800° C. for about 135 minutes. Disassembly of the graphitecontainment fixture revealed that the fibers of the fabric layers of thefabric preform were coated and bonded together by a boron nitridecoating. This boron nitride coating had a thickness of about 0.3 micronas determined by the weight gain of the fabric preform due to boronnitride deposition.

Next, a silicon carbide coating was then applied on top of the boronnitride coated and bonded fabric preform as follows. The boron nitridecoated and bonded fabric preform was suspended about 11.5 inches (282mm) from the top of the reactor using KANTHAL™ iron-chromium-aluminumalloy wires. The orientation of the boron nitride coated and bondedfabric preform was such that the warp yarns of the eighth layer of the12-harness satin weave fabric were parallel to the gas flow directionwithin the chamber. The rest of the silicon carbide coating process wassubstantially the same as that described in Example 10 for the finalsilicon carbide coating. The boron nitride bonded fabric preform wasfound to have been coated with a layer of silicon carbide, therebyforming a silicon carbide (SiC)/boron nitride (BN)-coated fabricpreform. The silicon carbide coating had a thickness of about 2.3microns as determined by the weight gain of the preform during siliconcarbide deposition.

A ceramic matrix comprising aluminum oxide and some aluminum alloy wasthen grown into the coated fabric preform in substantially the samemanner as was described for Sample A of Example 3.

Unlike the composite material described in Example 2, the compositematerial of the present example was neither subjected to the metalremoval process nor the elevated temperature heat treatment.

Sample O

The Sample O composite material was prepared in substantially the sameway as described for the Sample N composite material with the exceptionthat the coating temperature of about 800° C. for boron nitridedeposition was maintained for about 180 minutes. The resulting coatingthicknesses for boron nitride and silicon carbide were about 0.41 andabout 2.30 microns, respectively.

Sample P

The composite material of Sample P was produced in substantially thesame manner as was the material for Sample N with the exception that thecoating temperature of about 800° C. for deposition of boron nitride wasmaintained for about 70 minutes. The coating thicknesses of boronnitride and silicon carbide which resulted were about 0.20 and about2.30 microns, respectively.

Sample Q

The composite material of Sample Q was produced in substantially thesame manner as was described for Sample N with the exception that thetemperature of about 925° C. for deposition of the silicon carbidecoating was maintained for about 16 hours instead of about 20 hours. Thecoating thicknesses of boron nitride and silicon carbide which resultedwere about 0.29 and about 1.78 microns, respectively.

Sample R

The composite material of Sample R was produced in substantially thesame manner as was described for Sample Q with the exception that thetemperature of about 800° C. for deposition of the boron nitride coatingwas maintained for about 70 minutes. The coating thicknesses of boronnitride and silicon carbide which resulted were about 0.21 and about1.90 microns, respectively.

Test specimens for determining the mechanical strength of theabove-described composite materials were diamond machined to thespecimen dimensions given in Example 2.

The mechanical strength of these test specimens was then determined bystressing each specimen in four-point flexure substantially as describedin Example 2 until the sample failed. These flexural strength tests wereconducted at ambient temperature and are reported in Table V. Anexamination of Table V reveals that for a relatively constant siliconcarbide coating thickness of about 2.2 to 2.3 microns, the greatestflexural strength (taken from an average of 7 specimens for eachcomposite material sample) is realized for a boron nitride coatingthickness of about 0.3 micron. Reductions in strength were observed forboron nitride coating thicknesses of 0.20 and 0.41 microns,respectively. In addition, specifically by comparing the averagestrength values for Sample Q to that of Sample N and that of Sample R tothat of Sample P, it is clear that silicon carbide coating thicknessesin the range of 2.2-2.3 microns produce greater strengths than siliconcarbide coating thicknesses in the range of 1.8-1.9 microns.

Thus, this Example demonstrates that there exists a desirable range ofcoating thicknesses of each of boron nitride and silicon carbide to beapplied to the fibers of the present reinforced ceramic compositematerials to yield optimum ambient temperature strength. Specifically,the applied thickness of boron nitride should be greater than about 0.20micron but less than about 0.41 micron. Furthermore, the appliedthickness of silicon carbide should be greater than about 1.78 micronsand preferably greater than about 1.90 microns.

TABLE V Sample Vol. % BN* SiC* Avg. Str.** No. Fiber (μm) (μm) (MPa) N35.3 0.30 2.20 541 O 36.0 0.41 2.30 487 P 36.0 0.20 2.30 457 Q 36.0 0.291.78 388 R 36.0 0.21 1.90 358 *nominal coating thickness calculated onthe basis of weight gain during CVI **average of 7 specimens

EXAMPLE 14

This Example demonstrates the effect of the chemical-vapor-infiltration(CVI) coating thicknesses of both boron nitride and silicon carbide aswell as the volume fraction of reinforcement fibers on an elevatedtemperature flexural strength of a fiber reinforced ceramic compositematerial.

Samples S-X

Fabric preforms comprising ceramic grade NICALON® silicon carbide fiber(obtained from Dow Corning Corp., Midland, Mich.) was assembled insubstantially the same manner as was shown in Example 2. Unlike those ofExample 2, however, the fabric preform of the present Example onlymeasured about 3.2 inches (77 mm) square.

The fabric preforms were then chemical-vapor-infiltrated (CVI) withboron nitride (BN) and silicon carbide (SiC) in substantially the samemanner as was described in Example 2 with the following notableexceptions.

Twelve fabric preforms were simultaneously coated with boron nitride,and for the silicon carbide deposition, six boron nitride coatedpreforms were simultaneously coated. Furthermore, during the boronnitride deposition, a coating temperature of about 742° C. as indicatedby a thermocouple contained within the reactor chamber was maintainedfor about 5.5 hours and the total pressure within the reactor chamberwas maintained between about 1.2 and about 1.6 torr. The gaseousreactants consisted of ammonia (NH₃) flowing at a rate of about 1200standard cubic centimeters per minute (sccm) and boron trichloride(BCl₃) flowing at a rate of about 300 sccm. For the silicon carbidedeposition, a coating temperature of about 928° C. was maintained forabout 24 hours at a total reactor chamber operating pressure of about 11torr.

The differences in the boron nitride and silicon carbide coatingthicknesses which were obtained, as shown in Table VI, are accounted forby the relative location of a particular preform within the reactor.Generally speaking, for both boron nitride and silicon carbidedepositions, the closer the preform was to the gas reactant source, thethicker was the coating. During boron nitride deposition, the preformswere arranged six deep in groups of two. Likewise, during siliconcarbide deposition, the boron nitride coated preforms were arrangedthree deep in groups of two. Thus, during boron nitride deposition, theSample W preform was closest to the gas reactant source, while theSample V preform was farthest away. Likewise, during silicon carbidedeposition, the Sample X preform was closest to the gas reactant source,while the Sample T preform was farthest away.

TABLE VI Sample Vol. % BN* SiC* Avg. Str.** No. Fiber (μm) (μm) (MPa) S41 0.33 2.03 427 T 41 0.33 1.75 358 U 37 0.48 2.10 406 V 37 0.24 2.07377 W 33 0.51 2.20 367 X 32 0.49 2.43 312 *nominal coating thicknesscalculated on the basis of weight gain during CVI **average of 4specimens

Each fabric preform coated with boron nitride and silicon carbide wasthen infiltrated with a ceramic matrix comprising aluminum oxide andsome aluminum alloy using a directed metal oxidation processsubstantially as described in Example 2.

At least some of the metallic component of the formed fiber reinforcedceramic composite bodies was then removed. This metal removal processwas substantially as described in Example 2 with the followingexceptions. A nitrogen gas flow rate of about 10,000 sccm instead ofabout 4000 sccm was maintained throughout the heating and cooling. Also,three fiber reinforced composite bodies were simultaneously processed.

TABLE VII Matrix Fabrication Parameters Residual Metal Example Sample IDFiber Plies Atmosphere Temperature Time (h) Removed Comments 1 — 3bundles of 500 fibers O₂ 1000° C.  54 no 2 — 12 HSW-(6) 8 HSW-12 HSW air950° C. 90 yes Some heat treated (see Table I); parent metal containing8.5-11 wt % Si 3 A (8) 12 HSW air 1000° C.  60 yes 3 B (4) 12 HSW N₂900° C. 200 yes 4 C 12 HSW-(6) 8 HSW-12 HSW air 950° C. 90 no 4 D 12HSW-(6) 8 HSW-12 HSW air 950° C. 90 yes 4 E 12 HSW-(6) 8 HSW-12 HSW air950° C. 90 yes 4 F 12 HSW-(6) 8 HSW-12 HSW air 950° C. 90 yes 5 — 12HSW-(6) 8 HSW-12 HSW air 950° C. 90 yes 6 — 12 HSW-(6) 8 HSW-12 HSW air950° C. 90 yes 7 — 12 HSW-(6) 8 HSW-12 HSW air 950° C. 90 yes 8 — (2) 8HSW-(3) 12 HSW-(2) 8 HSW air 950° C. 72 yes 8 — (2) 8 HSW-(3) 12 HSW-(2)8 HSW air 950° C. 100 yes 9 G, H, I (8) 12 HSW air 950° C. 90 yes 10 —(8) 12 HSW air 1000° C.  60 yes 12 — (8) 8 HSW air 950° C. 90 yes 5layers of SiC particulate 13 N-R (8) 12 HSW air 1000° C.  60 no 14 S-X12 HSW-(6) 8 HSW-12 HSW air 950° C. 90 yes

Unlike the method of Example 2, however, no subsequent heat treatment ofthe formed ceramic composite bodies of the present Example wasperformed.

The mechanical strength of each of the fiber reinforced ceramiccomposite bodies of Samples S-X was measured at a temperature of about1200° C. in a four-point flexure mode. Specifically, test specimens werediamond machined to the dimensions specified in Example 2 and tested infour-point bending until failure using the procedure as outlined inExample 2. The resulting flexural strength based upon an average of fourdata points per material is reported for each of Samples S-X in TableVI.

An examination of the data reported in Table VI reveals a number oftrends.

Specifically, a comparison of the data for Sample S with those forSamples U, V and W shows that boron nitride coating thicknesses of about0.24 micron and 0.48 micron or thicker resulted in flexural strengthswhich were suboptimal. Thus, for the fiber reinforced composite bodiesof the present invention, a boron nitride coating thickness on thereinforcement fibers of about 0.3 micron ought to be at least close tooptimal.

Moreover, inspection of Table VI reveals a correlation between thethickness of the boron nitride coating and the volume fraction of theNICALON® reinforcement fiber in the coated preform (and ultimately thecomposite body). This correlation results from the “spring-back”phenomenon which occurs after coating the fabric preforms with boronnitride when the graphite containment fixture is disassembled and theboron nitride coated fabric preform is removed. It seems that thethicker the boron nitride coating which is applied, the more the coatedfabric preform expands in its thickness dimension upon its removal fromthe graphite containment fixture, thus the lower the overall volumefraction of NICALON® fiber.

In addition, a comparison of the strength data for Samples S, U and Wshows that the flexural strength of the composite material increases asthe volume fraction of reinforcement fibers increases.

Finally, a comparison of the data for Sample T with the data for SampleS and, similarly, a comparison of the data for Sample X with the datafor Sample W reveals that for a given volume fraction of reinforcementfibers, a silicon carbide thickness of about 1.75 microns, while athickness of about 2.43 microns of silicon carbide is excessive in thatit also results in suboptimal strength. The importance of the propersilicon carbide thickness is highlighted by a comparison of the data forSample T with for Sample W. Specifically, this comparison shows that theoptimization of the silicon carbide coating thickness at about 2.2microns compensates for a difference in volume fraction of reinforcementfibers of about eight points or about 20 percent.

Thus, this Example demonstrates that for the fiber reinforced ceramiccomposite materials described herein, there exists desirable boronnitride and silicon carbide coating thicknesses to apply to thereinforcement fibers to optimize the high temperature strength of theformed composite materials. Specifically, the thickness of boron nitrideapplied should be greater than about 0.24 micron but less than about0.51 micron. The thickness of silicon carbide applied should be greaterthan about 1.75 microns but less than about 2.43 microns. These desiredthickness ranges agree well with the optional ranges of Example 13derived for optional ambient temperatures composite strength.Furthermore, within the range of fiber reinforcement of about 32 to 41volume percent of the composite body, the strength of the composite bodyincreases with the volume fraction of reinforcement.

Table VII summarizes many of the processing parameters followed infabricating the fiber reinforced ceramic composite bodies described inthe foregoing Examples.

EXAMPLE 15

This Example demonstrates, among other things, an improved method ofcoating a fabric preform. Specifically, this Example demonstrates a setof coating conditions which result in coatings of more uniform thicknessthroughout the fabric preform.

A fabric preform 103 was made by stacking a plurality of layers of 8harness satin weave (8 HSW) fabric and 12 harness satin weave (12 HSW)fabric made from ceramic grade NICALON® silicon carbide fiber (obtainedfrom Dow Corning Corp., Midland, Mich.) on top of each other insubstantially the same manner as was described in Example 8. The fabricpreform had dimensions of about 9 inches (229 mm) long by about 6 inches(152 mm) wide by about 0.125 inch (3.2 mm) thick.

The fabric preform was clamped in substantially the same kind of fixtureas was described in Example 2 and depicted in FIG. 5e. The preformcontainment fixture 108 containing the fabric preform was placed into areactor chamber of a refractory alloy steel chemical vapor infiltration(CVI) apparatus having a graphite tube liner and having overalldimensions of about 8 feet (2.4 meters) in length by about 15.5 inches(394 mm) in inside diameter. The warp yarns of the first and seventhlayers of the fabric preform were perpendicular to the gas flowdirection within the chamber as well as being perpendicular to thelongitudinal axis of the reactor chamber. The reactor chamber was closedand evacuated to less than about 0.5 torr. The reactor chamber was thenheated to a temperature of about 800° C. When the temperature within thereactor TABLE VII chamber reached about 800° C., a gas mixturecomprising borontrichloride (BCl₃) flowing at about 1200 sccm at atemperature of about 60° C. and ammonia (NH₃) flowing at about 2100 scmwas introduced into the reactor chamber while maintaining a totaloperating pressure of about 0.5 torr. After about 4 hours at atemperature of about 800° C., the gas mixture flowing into the reactorchamber was interrupted, the power to the furnace heating the reactorchamber was interrupted and the furnace and its contents were naturallycooled. At a temperature below about 200° C., the reactor chamber doorwas opened and the graphite containment fixture was removed, cooled anddisassembled to reveal that the fibers of the fabric layers of thefabric preform were coated and that the fabric layers comprising thefabric preform were bonded together by a boron nitride coating. Theboron nitride coating had a thickness of about 0.48 micron.

The boron nitride coated fabric preform was then stored in a vacuumdesiccator until it was ready to be put back into the chemical vaporinfiltration apparatus for additional coating.

For the application of this subsequent coating, the boron nitride coatedand bonded fabric preform was placed back into the reactor chamber ofthe chemical vapor infiltration apparatus. In this instance, however,the warp yarns of the first and seventh layers of the fabric preformwere parallel to the gas flow direction within the chamber, as well asbeing parallel to the longitudinal axis of the reactor chamber. Morespecifically, the boron nitride coated fabric preforms were supported bya graphite fixture as shown in FIG. 18A. The graphite fixture alone isshown in FIG. 18B. A total of 8 boron nitride coated fabric preforms canbe further coated simultaneously in a single reactor run by placing 2such loaded fixtures front-to-back in the reactor chamber.

The CVI reactor chamber was closed and evacuated to about less thanabout 1 torr. Hydrogen gas was introduced into the reactor chamber at aflow rate of about 11,000 standard cubic centimeters per minute (sccm).The reactor chamber was then heated to a temperature of about 950° C.The reactor pressure was equilibrated at about 250 torr. Once thetemperature of the contents of the reactor chamber had substantiallycompletely stabilized at about 950° C., about 180° sccm of hydrogen werediverted away from direct entry into the reactor chamber and were firstbubbled through a bath of methyltrichlorosilane (MTS) maintained at atemperature of about 45° C. before entering the reactor chamber. Afterabout 48 hours at a temperature of about 950° C., the power to thefurnace heating the reactor chamber was interrupted and the about 180°sccm of hydrogen that was being directed through the MTS bath was againpermitted to flow directly into the reactor chamber to re-establish adirect hydrogen gas flow rate of about 11000 sccm into the reactorchamber. After the reactor chamber had cooled substantially, thehydrogen flow rate was interrupted and the furnace and its contents wereevacuated to less than 1 torr. The pressure within the reactor chamberwas then brought back up to about atmospheric pressure with argon gas.After the reactor chamber had cooled to a temperature below about 200°C., the argon gas flow rate was interrupted and the reactor chamber doorwas opened. The graphite support fixtures were removed, cooled anddisassembled to reveal that the boron nitride bonded fabric preforms hadbeen coated with a second layer of silicon carbide thereby forming asilicon carbide (SiC)/boron nitride (BN)-coated fabric preform. Thesilicon carbide had a thickness of about 2-3 microns. Significantly, thesilicon carbide coating was of more uniform thickness from the interiorof the preform to an exterior surface in the present Example than in thepreviously described Examples. In other words, the thickness of siliconcarbide deposited at the exterior of the preform was not as great as inearlier Examples; thus, the coated preforms of the present Example weremore permeable than some of the coated preforms of the previousExamples. Thus, the results of this Example suggest that it may bepossible to apply silicon carbide coatings to the present fabricpreforms having nominal thickness greater than about 2 to 3 micronswithout creating isolated pores in the preform.

Growth of an alumina oxidation reaction product through the siliconcarbide/boron nitride-coated fabric preform was then carried out insubstantially the same manner as was described in Example 2 to form afiber reinforced ceramic composite body comprising a ceramic matrixcomprising an aluminum oxide oxidation reaction product and a metalliccomponent comprising some residual unreacted parent metal, with theceramic matrix embedding the silicon carbide/boron nitride coatedNICALON® silicon carbide fibers. Substantially complete growth of theceramic matrix only required about 72 hours, however. Because of themore permeable nature of the coated fabric preforms of the presentExample, it is believed that the time required for complete growth iseven less than this value.

Thus, this Example demonstrates an efficient technique for coating aplurality of preforms simultaneously as well as conditions which resultin silicon carbide coatings of more uniform thickness.

EXAMPLE 16

This Example demonstrates the fatigue characteristics of the presentfiber reinforced ceramic composite materials. Specifically, this Exampledemonstrates the lifetimes for samples of ceramic grade NICALON® siliconcarbide fiber reinforced alumina matrix composite bodies as a functionof the maximum applied stress for bodies tested in air at varioustemperatures and subjected to low-frequency cycling in tension.

Samples which were tested at about 20° C. and at about 1000° C. werefabricated in substantially the same manner as was described in Example2, including the residual metal removal process. The fiber reinforcedceramic composite bodies which were tested at temperatures of about1100° C. and about 1370° C. were fabricated substantially in accordancewith Example 8, which fabrication also included the residual metalremoval process.

The geometry of the test specimen was that of a “double dogbone”, thatis, similar to that geometry in FIG. 7 except further comprising anotherreduced section. Specifically, test specimens were diamond machined toan overall length of about 5 inches (127 mm), about 0.55 inch (14 mm)maximum width and having a gage section measuring about 1.3 inches (33mm) in length by about 0.25 inch (6 mm) in width.

Each sample was placed into the test chamber of a universal testingmachine at about 25° C. The test chamber was then heated to the desiredelevated temperature in air. When the temperature of the specimen hadstabilized, a sinusoidal tensile stress was applied to the specimen. Theminimum applied tensile stress was about 10 percent of the maximumapplied tensile stress. The testing apparatus was configured so as torecord the number of tensile stress cycles required to cause failure.These fatigue data are illustrated in Table VIII. Those test specimenswhich were still intact following 10,000 cycles of applied tensilestress at temperatures of 1100° C. or 1370° C. were then tensile testedin air using a uniformly increasing load at the same temperature atwhich they were cycled in applied stress until failure was observed. Thetable shows that such test specimens retained over 50 percent of theiroriginal strength following the tensile cycling at elevated temperature.The table also demonstrates that the fiber reinforced ceramic compositematerial was capable of surviving over 2 million cycles of a tensilestress applied between about 8 and about 83 MPa at a frequency of about5 Hz at a temperature of about 1000° C. in air. The data generated atthe elevated temperatures are presented graphically in FIG. 19 whichshows the number of cycles to produce failure as a function of themaximum applied tensile stress. The arrows connected to several of thedata points and pointing to the right indicate that the particular datapoint represents a lower bound of the material's life (e.g., failure ofthe specimen had not been achieved at the indicated maximum stress andnumber of tensile cycles). Further examination of the data presented inFIG. 19 suggest an endurance limit for the fiber reinforced ceramiccomposite material of about 80 MPa at a temperature of about 1000° C. inair.

TABLE VIII 2-D Nicalon ™ /Al₂O₃ Low Cycle Fatigue Data Max. StressResidual Str. Test (R = 0.1) Frequency Cycles at Temp. Temperature MPa(ksi) Hz to failure MPa (ksi) Room Temperature 72 (11) 1 172,912* — 72(11) 1 235,858* — 138 (20) 1 150,956* — 138 (20) 1 166,459* — 166 (24)  1/10 351,600/118,360 — 1000° C. (1800° F.) 83 (12) 5 2,195,000* — 103(15)   0.3 142,577* — 103 (15) 5 117,843 — 103 (15) 5 175,620 — 124 (18)1 27,712 — 124 (18) 5 66,676 — 1100° C. (2000° F.) 60 (9) 1 10,000* 185(26) 120 (17) 1 10,000* 149 (21) 180 (26) 1 603 — 1370° C. (2500° F.) 58(8) 1 10,000* 153 (22) 115 (16) 1 2,480 — 173 (25) 1 128 — *Test stoppedprior to failure RT & 1000° C. data generated by GE 1100 & 1370° C. datagenerated by Williams International

EXAMPLE 17

This Example illustrates a modified stress rupture test whose purpose orobjective was to further simulate at least some of, the conditions whichmight be present in a turbine engine. Specifically, the test illustratedby the present example is similar in many respects to the stress rupturetest described in Example 9 with the exception that a temperature cyclewas added or superimposed to the test system during the elevatedtemperature exposure under the applied dead load.

A fiber reinforced ceramic composite body was fabricated insubstantially the same manner as was described in Example 2, includingremoving a substantial fraction of the residual metallic component.

Tensile test specimens were diamond machined in substantially the samemanner as was described in Examples 7 and 9 and loaded into the testfixture described in Example 9, which in turn was loaded into the ModelP-5 creep testing machine (SATEC Inc., Grove City, Pa.). A tensilestress of about 12.5 MPa was then applied to the test specimen usingdead loading. A resistance-heated air atmosphere furnace was positionedcompletely around the test fixture portion of the creep testing machineand the furnace and the sample contained within were heated from about20° C. to a temperature of about 1100° C. Each thermal cycle thenconsisted of maintaining a temperature of about 1100° C. for about 1hour in air, then uniformly decreasing the temperature of the testspecimen to a temperature of about 600° C. over a period of about 45minutes, maintaining a temperature of about 600° C. for about 1 hour andfinally uniformly increasing the temperature of the specimen back up toa temperature of about 1100° C. over a period of about 45 minutes. As inExample 9, sample strain was monitored with a Model 1102 ZYGOhelium-neon laser extensometer (Zygo Corporation, Middlefield, Conn.).The sample test data were then recorded in the form of sample strain asa function of test duration, which data are illustrated graphically inFIG. 20. Referring to FIG. 20, after the application of about 12 thermalcycles at an applied tensile stress of about 25 MPa (the first fewcycles being used to check out the functioning of all of the testequipment), the stress was then increased to about 50 MPa. After theapplication of about 115 thermal cycles at a stress of about 50 MPa, thestress was then further increased to about 70 MPa. After about 22 cyclesat a stress of about 70 MPa, the sample failed. FIG. 20 also shows anenlargement of the 70 MPa region which specifically illustrates thechange in strain in the test specimen in response to the change inspecimen temperature. All together, the test specimen survived a totalof 142 thermal cycles or about 500 hours of test duration.

A second thermal cycling test was then conducted on a substantiallysimilar fiber reinforced composite test specimen in substantially thesame manner as described above with the exception that the appliedtensile stress throughout the duration of the testing was about 50 MPa.This second test specimen survived thermal cycling under this appliedtensile stress in air for a total of about 282 thermal cycles (or about987 hours) before failure occurred.

Thus, this Example demonstrates that the present fiber reinforcedceramic composite materials are capable of surviving hundreds of hoursunder tensile loads in air under conditions of varying elevatedtemperature.

EXAMPLE 18

This Example demonstrates, among other things, the deposition of achemically modified coating layer on a reinforcement filler material andsubsequent encapsulation by a matrix material to form a composite body.More specifically, the present Example demonstrates the incorporation ofa source of silicon into a boron nitride based coating material.

SAMPLE AA (no silicon doping)

A fabric preform comprising ceramic grade Nicalone® silicon carbidefiber was fabricated in substantially the same manner as described inExample 2. The fabric preform of the present Example measured about 3inches (76 mm) square by about 0.125 inch (3 mm) thick.

The fabric preform was then coated with a material comprising boronnitride in substantially the same manner as in Example 2 except that theusable inside diameter of the coating chamber was about 5 inches (127mm) instead of about 9.45 inches (240 mm), the temperature inside thereactor was maintained at about 736-740° C., the operating pressure wasmaintained at about 1.1 to 1.2 Torr, the time at temperature was about 4hours, the flowrate of the ammonia reactant was about 342 standard cubiccentimeters per minute (sccm), and the flowrate of the boron trichloridewas about 85 sccm. An average of about 0.32 micron of boron nitride wasdeposited on the Nicalon® silicon carbide filaments, as calculated fromthe weight gain of the fabric preform.

Next, a coating comprising silicon carbide was deposited on top of theboron nitide coated fabric preform. The deposition conditions weresubstantially the same as those described for the silicon carbidedeposition of Example 2 except that the coating was deposited at atemperature of about 980° C., at a pressure of about 11 Torr and and fora duration of about 19 hours. Based upon the preform weight gain, it wasestimated that the average silicon carbide coating thickness on thefibers was about 1.32 microns.

Following this CVI) coating process, a matrix comprising aluminum oxidewas formed by directed metal oxidation of a parent metal comprisingaluminum. More specifically, the matrix was formed in substantially thesame manner as the matrix described in Example 2 with the exception thatthe nickel oxide particulate was about minus 200 mesh (substantially allparticles smaller than about 75 microns) and the dwell temperature ofabout 950° C. was maintained for about 96 hours.

Following matrix formation, most of the residual, unreacted parent metalwithin the formed alumina matrix composite was removed. The metalremoval technique was substantially the same as that described inExample 2 except that the metal removal lay-up was loosely covered witha Grafoil® graphite foil lid (Union Carbide Co., Cleveland, Ohio), andthat the filler material mixture for infiltration comprised by weightabout 10 percent ground magnesium particulate (−100 +200 mesh, HartMetals, Tamaqua, Pa.) and the balance grade C75 unground aluminaparticulate (Alcan Chemicals Div. of Alcan Aluminum Corp., Cleveland,Ohio). A weight loss in the composite body of about 8.4 percent wasrecorded.

Unlike some of the flexural test specimens in Example 2, the compositebody of the present Example was not heat treated.

SAMPLE AB

Sample AB features an attempt to deliberately add a source of silicon tothe CVD gas stream for the purposes of doping the resulting boronnitride based coated with silicon.

Sample AB was fabricated by substantially the same procedures as wasSample AA with the following exceptions: The boron nitride coating runwas conducted at a pressure of about 1.6 to 1.7 Torr. Also, to the NH₃and BCl₃ gas streams was added about 200 sccm of hydrogen (H₂) andsilicon tetrachloride (SiCl₄). The H₂ was bubbled through liquid SiCl₄at ambient (e.g., about 20° C.) temperature, thereby acting as a carriergas for SiCl₄. About 0.32 micron of the BN based coating was applied tothe fibers of NICALON® silicon carbide, based upon the weight gain ofthe fabric preform.

During the silicon carbide deposition by CVD, the dwell at about 980° C.was maintained for about 11.5 hours. About 1.79 microns of siliconcarbide were deposited.

SAMPLE AC

Sample AC was prepared in substantially the same manner as was Sample ABwith the following exceptions: The boron nitride based coating wasdeposited for about 5 hours at a temperature of about 735° C. to 740° C.and at a pressure of 1.3-1.4 Torr. The BCl₃ and NH₃ gas flowrates wereabout 65 sccm and 262 sccm, respectively. The SiCl₄ with its H₂ carrierwas admitted to the coater at a flowrate of about 496 sccm. A coatingabout 0.42 micron thick was deposited. The silicon carbide coating wasdeposited identically to that described for Sample AA.

SAMPLE AD

Sample AD is also a silicon carbide fiber reinforced alumina matrixcomposite material wherein the fibers are coated with a duplex coatingfeaturing a boron nitride containing layer followed by a silicon carbidecoating layer. As with Samples AB and AC, the BN layer of the presentsample was also modified, but by a different route. Instead of the SiCl₄being carried into the reactor by H₂, however, pure SiCl₄ gas with nocarrier was employed. In particular, it was realized that the vaporpressure of SiCl₄ at ambient temperature was sufficient to produce aflowrate of about 21 sccm under the process conditions. The BCl₃ and NH₃gas flowrates were 85 sccm and 342 sccm, respectively. The modified BNlayer was deposited for about 5 hours at a pressure of about 1.3 to 1.4Torr. An approximately 0.37 micron thick coating was deposited.

The SiC deposition was identical to that in Sample AA.

SAMPLE AE

Sample AE was prepared almost identically to Sample AD. The onlysignificant difference between these two Samples is that for the BNbased coating, the present Sample featured gas flowrates of 60 sccm, 69sccm and 276 sccm for the SiCl₄, BCl₃ and NH₃, respectively. Anapproximtely 0.35 micron thick coating was deposited.

SAMPLE AF

Sample AF was prepared in substantially the same manner as was Sample AEexcept that the gas flowrates for the modified BN deposition were 40sccm, 82 sccm and 326 sccm of SiCl₄, BCl₃ and NH₃, respectively. Thethickness of the resulting coating was about 0.35 micron.

A chemical elemental analysis was performed on the modified BN coatingson some of the Samples. This elemental analysis was performed by anoutside contractor using scanning Auger depth profiling. In addition tothe expected B, N and Si, considerable quantities of C and O were alsodetected. The results of this elemental analysis are reported in termsof atomic percent in Table IX. The column labeled “Si:B ratio” providesthe ratio of silicon to boron atoms in the precursor SiCl₄ and BCl₃gases.

To guage the effect of the attempted silicon modification of the boronnitride coating on the strength of the composite bodies into which thecoated were encapsulated, a number of flexural strength sample test barswere diamond machined from each Sample. The four point flexural strengthtesting was performed at ambient temperature and at about 1200° C. inair substantially as described in Example 2. The mean flexural strengthbased on a sample size of 3 is reported in Table X, as well as presentedin graphical form in FIG. 21. The flexural strength shows littledependence upon the ratios of the boron and silicon precursor gases.

To assess the resistance of the present composite bodies to chemicalattack by oxygen and moisture at elevated temperatures, three of thepresent Samples were selected for such corrosion testing. Specifically,Samples AD, AE and AF were subjected to an atmosphere consisting of 90percent by volume water vapor, balance oxygen at a temperature of about800° C. and ambient pressure. After an approximately 88 hour exposure,the recession distance (e.g., corrosion length) was measured from amachined surface of the composite body. This recession distance wasabout 600 microns to 800 microns for Sample AE and about 500 microns toabout 700 microns for Sample AF, but only about 2 microns to 40 micronsfor Sample AD. By comparison, a typical Si/SiC matrix composite bodyfeaturing an unmodified boron nitride fiber coating exhibits recessiondistances of about 1700 microns to about 2200 microns. Although thematrix material is not thought to significantly affect the recessionrate, since all test samples feature exposed filament ends, if anything,the alumina matrix composites of Sample AD, AE and AF might be expectedto corrode faster than the composite body formed by melt infiltrationsince the alumina matrix has a higher population density of microcracks.The fact that these Samples exhibited less corrosion (less recessiondistance) suggests that the modified BN coating on the fibers is morechemically protective than regular , unmodified BN.

An artisan of ordinary skill will readily appreciate that numerousmodifications may be made to the above-identified Examples withoutdeparting from the spirit of the present invention. Accordingly, theExamples should be considered as illustrative of the invention andshould in no way be construed as limiting the scope of the invention asdefined in the claims appended hereto.

TABLE IX Si:B Atom Elemental Analysis, Atom % Sample ratio (precursorgases) B N Si C O AA 0 AB 2.3 AC 7.7 39 38 0 18 5 AD 0.25 39 41 1-2 13 5AE 0.87 37 42 3 10 10  AF 0.49 38 40 3 13 6

TABLE X Si:B Atom Four Point Flexural Strength (MPa) Sample ratio(precursor gases) at 20° C. at 1200° C. AA 0 483 +/− 51 374 +/− 15 AB2.3 328 +/− 29 251 +/− 5  AC 7.7 455 +/− 16 311 +/− 31 AD 0.25 429 +/−20 307 +/− 21 AE 0.87 360 +/− 45 320 +/− 36 AF 0.49 427 +/− 53 362 +/−10

What is claimed is:
 1. A composite material, comprising: (a) a permeablemass or preform comprising a plurality of bodies of at least one fillermaterial; (b) a matrix embedding said permeable mass or preform; and (c)at least one coating disposed between said bodies of at least one fillermaterial and said matrix, said at least one coating comprising boronnitride having at least some silicon present in solution in said boronnitride; and (d) at least two zonal junctions, at least one of saidzonal junctions being weak relative to the remaining zonal junction(s)to permit debonding and pull-out of said at least one filler materialwith respect to said matrix upon application of stress sufficient tocause fracture of said composite material.
 2. The composite material ofclaim 1, further comprising at least one oxide glass network-formerdisposed between said bodies of at said at least one filler material andsaid matrix.
 3. The composite material of claim 2, wherein said at leastone oxide glass network-former is provided as a particulate slurry. 4.The composite material of claim 2, wherein said at least one oxide glassnetwork-former comprises silica.
 5. The composite material of claim 1,wherein said matrix comprises a material selected from the groupconsisting of silicon, silicon carbide and aluminum oxide.
 6. Thecomposite material of claim 1, further comprising at least one oxygengetterer.
 7. The composite material of claim 1, wherein said debondingoccurs at an interface between a coating and (a) said matrix, (b) saidfiller material or (c) another coating, and not within a coating.
 8. Thecomposite material of claim 1, wherein said matrix is produced by amethod selected from the group consisting of directed metal oxidationand melt infiltration.
 9. The composite material of claim 1, whereinsaid filler material comprises silicon carbide.
 10. The compositematerial of claim 1, wherein at least one of said plurality of bodiescomprises a filament or fiber, and further wherein said at least onecoating is coextensive with a longitudinal axis of said fiber.
 11. Thecomposite material of claim 1, further comprising at least oneprotective coating disposed between said matrix and said coating thatcomprises said boron nitride.
 12. The composite material of claim 11,wherein said protective coating comprises at least one material selectedfrom the group consisting of silicon carbide, silicon nitride andaluminum oxide.
 13. The composite material of claim 1, wherein said atleast one coating is at least partially amorphous.
 14. The compositematerial of claim 1, wherein said at least one coating exhibits limitedcrystallinity.
 15. The composite material of claim 1, wherein said atleast one coating comprises a plurality of regions or “domains” eachabout 5 to about 20 nanometers in size, and wherein said regions ordomains exhibit a lamellar crystal structure.
 16. The composite materialof claim 15, wherein a lamellar crystal structure within a given regionor domain is randomized in spatial orientation with respect to alamellar structure in a different region or domain.
 17. The compositematerial of claim 1, wherein said boron nitride comprises at least about0.5 atomic percent silicon.
 18. The composite material of claim 1,wherein said boron nitride comprises up to about 3 atomic percentsilicon.
 19. The composite material of claim 1, wherein said boronnitride comprises from about 1 to about 3 atomic percent silicon. 20.The composite material of claim 1, wherein said at least one coatingfurther comprises oxygen.
 21. The composite material of claim 1, whereinsaid at least one coating further comprises carbon.
 22. A compositematerial, comprising: (a) a permeable mass or preform comprising aplurality of bodies of at least one filler material; (b) a matrixembedding said permeable mass or preform; and (c) at least one coatingdisposed between said bodies of at least one filler material and saidmatrix, said at least one coating comprising boron, nitrogen and no morethan about 3 atomic percent silicon; and (d) at least two zonaljunctions, at least one of said zonal junctions being weak relative tothe remaining zonal junction(s) to permit debonding and pull-out of saidat least one filler material with respect to said matrix uponapplication of stress sufficient to cause fracture of said compositematerial.
 23. The composite material of claim 22, wherein said at leastone coating comprises a boron nitride structure.
 24. The compositematerial of claim 22, wherein said at least one coating comprises boronnitride modified by said silicon.
 25. A composite material, comprising:(a) a permeable mass or preform comprising a plurality of bodies of atleast one filler material; (b) a matrix embedding said permeable mass orpreform; and (c) at least one coating disposed between said bodies of atleast one filler material and said matrix, said at least one coatingcomprising boron and nitrogen in roughly equal atomic proportions and nomore than about 3 atomic percent silicon; and (d) at least two zonaljunctions, at least one of said zonal junctions being weak relative tothe remaining zonal junction(s) to permit debonding and pull-out of saidat least one filler material with respect to said matrix uponapplication of stress sufficient to cause fracture of said compositematerial.
 26. A composite material, comprising: (a) a permeable mass orpreform comprising a plurality of bodies of at least one fillermaterial; (b) a matrix embedding said permeable mass or preform; and (c)at least one coating disposed between said bodies of at least one fillermaterial and said matrix, said at least one coating comprising boronnitride modified by silicon atoms, said silicon making up no more thanabout 3 atomic percent of said boron nitride; and (d) at least two zonaljunctions, at least one of said zonal junctions being weak relative tothe remaining zonal junction(s) to permit debonding and pull-out of saidat least one filler material with respect to said matrix uponapplication of stress sufficient to cause fracture of said compositematerial.
 27. A composite material, comprising: (a) a permeable mass orpre-form comprising a plurality of bodies of at least one fillermaterial; (b) a matrix embedding said permeable mass or pre-form; (c) atleast one coating disposed between said bodies of at least one fillermaterial and said matrix, said at least one coating comprising boronnitride and silicon co-deposited by a chemical vapor deposition processat a temperature no greater than about 800° C.; and (d) at least twozonal junctions, at least one of said zonal junctions being weakrelative to the remaining zonal junction(s) to permit debonding andpull-out of said at least one filler material upon application of stresssufficient to cause fracture of said composite material.